Thesis Ndlovu

The Wear Properties of Tungsten Carbide-Cobalt Hardmetals

from the Nanoscale up to the Macroscopic Scale

Der Technischen Fakultät der Universität Erlangen-Nürnberg

zur Erlangung des Grades

DOKTOR-INGENIEUR

by

Siphilisiwe Ndlovu

Erlangen 2009

Das Verschleißverhalten von

Wolframkarbid-Kobalt-Hartmetallen von der Nanoskala bis

zur Makroskala

Der Technischen Fakultät der Universität Erlangen-Nürnberg

zur Erlangung des Grades

DOKTOR-INGENIEUR

vorgelegt von

Siphilisiwe Ndlovu

Erlangen 2009

Als Dissertation genehmigt von

Der Technischen Fakultät der

Universität Erlangen-Nürnberg

Tag der Einreichung: 22. 09. 2009

Tag der Promotion : 08. 12. 2009

Dekan: Prof. Dr.-Ing. habil. R. German

Berichterstatter: Prof. Dr. rer.nat. M. Göken

Prof. Dr. S. Virtanen

Abstract

A study has been conducted on the tribological properties of WC-Co hardmetals by

carrying out a series of wear tests from the nanoscale up to the macroscopic scale.

The composition of the hardmetals was varied and the binder content ranged from 6

to 15 wt%. The binder in all the samples was cobalt and one of the samples had small

additions of Cr3C2 and VC. The WC grain size in the samples ranged from 250 nm

(nano-size) up to 2.65 µm (coarse-grained). A binderless WC sample and a pure

cobalt sample were also included in this work and the mechanical properties of all

the samples were measured using nanoindentation. The wear tests on the nanoscale

consisted of scratch testing using a Nanoindenter XP. Macroscopic wear was

investigated by conducting three body abrasive and sliding wear tests.

The binderless WC sample was found to have the best performance at loads below

1 N. In this load range a smaller grain size led to an increase in the scratch resistance

for samples with 6 wt% binder. On the other hand for the samples with 15 wt%

binder, a smaller grain size resulted in a decrease in the scratch resistance. In the 1 to

10 N load range the binderless WC underwent brittle wear which led to very high

scratch depths. Whereas for the samples containing 15 wt% Co a smaller grain size

resulted in an increase in the scratch resistance. The finer grained hardmetals

exhibited lower abrasive wear rates than their coarse grained counterparts. The main

wear mechanisms were found to be plastic deformation via glide activity,

microcracking, binder extrusion and grain fall out.

Zusammenfassung

Die Untersuchung der tribologischen Eigenschaften von WC-Co-Hartmetallen von

der Nanoskala bis zur makroskopischen Skala stand im Vordergrund dieser Arbeit.

In der Zusammensetzung der Hartmetallproben wurde der Bindergehalt zwischen 6

und 15 Gew.% variiert. Als Binder wurde in allen Proben Kobalt verwendet; zudem

wurden bei einer Probe kleine Mengen von Cr3C2 und VC zugegeben. Die WC-

Korngröße der Proben lag zwischen 250 nm (nanokörnig) und 2,65 µm (grobkörnig).

Ebenso wurden eine binderfreie WC-Probe und eine reine Kobalt Probe untersucht.

Die mechanischen Eigenschaften und Verschleißmechanismen auf lokaler Ebene

wurden anhand von Nanoindentierungsexperimenten und Scratch-Versuchen

untersucht, die am Nanoindenter XP durchgeführt wurden. Zur Untersuchung des

Verschleißverhaltens auf der makroskopischen Skala wurden Dreikörper-

Abrasivverschleiß- und Gleitverschleißversuche herangezogen.

Die binderfreie WC-Probe zeigte das beste Verschleißverhalten bei Scratch-

Versuchen unter 1 N Last. Im Lastbereich von 5 bis 500 mN wurde für Proben mit

einem geringen Bindergehalt (6 Gew.%) für kleinere Korngrößen eine Erhöhung der

Kratzfestigkeit festgestellt. Im Gegensatz dazu führte bei Proben mit einem hohen

Bindergehalt (15 Gew.%) eine kleinere Korngröße zu einem Rückgang der

Kratzfestigkeit. Im Lastbereich von 1 bis 10 N wurde für die Proben mit 6 Gew.%

Kobalt-Binder keine Korngrößen-Abhängigkeit beobachtet. Jedoch in den Proben, die

15 Gew.% Binder enthalten, führte eine kleinere Korngröße zu einer Zunahme des

Kratzwiderstands. Die binderfreie WC-Probe zeigte eine Abnahme der

Kratzfestigkeit bei höheren Lasten. Bei den makroskopischen Versuchen führte eine

Verkleinerung der WC-Korngröße zu niedrigeren Verschleißraten. Plastische

Verformung der Wolframkarbid-Körner, Bildung von Mikrorissen, Binderabtrag und

Ausfall von WC-Körnen sind die vorliegenden Hauptverschleißmechanismen.

Table of Contents i

1 Introduction ..................................................................................................................... 1

2 Objectives ......................................................................................................................... 2

3 WC-Co hardmetals ......................................................................................................... 3 3.1 Introduction ........................................................................................................................ 3 3.2 Hard metal manufacture ................................................................................................... 3 3.3 The constituent phases...................................................................................................... 4 3.4 Mechanical properties of hard metals............................................................................ 8

4 Friction and Wear ......................................................................................................... 17 4.1 Friction ............................................................................................................................... 17 4.2 Introduction to Wear ....................................................................................................... 18

5 Wear of WC-Co hardmetals ........................................................................................ 24

5.1 Deformation behaviour of hardmetals ........................................................................ 24 5.2 Sliding wear of WC-Co hard metals............................................................................. 28 5.3 Scratch testing of WC-Co hard metals ......................................................................... 32 5.4 Abrasive wear of WC-Co hard metals.......................................................................... 35

6 Tribological testing....................................................................................................... 39

6.1 Sliding wear testing......................................................................................................... 39 6.2 Abrasive wear testing ...................................................................................................... 40 6.3 Scratch testing................................................................................................................... 41 6.4 Nanoindentation testing ................................................................................................. 44 6.5 Microstructure analysis................................................................................................... 49

7 Experimental methods .................................................................................................. 53

7.1 Materials and sample preparation ................................................................................ 53 7.2 Hardness and fracture toughness.................................................................................. 55 7.3 Nanoindentation .............................................................................................................. 56 7.4 Instrumented scratch testing.......................................................................................... 57 7.5 Three body abrasive wear tests ..................................................................................... 60 7.6 Sliding wear tests ............................................................................................................. 61

8 Results............................................................................................................................. 63

8.1 Material microstructure .................................................................................................. 63 8.2 Nanoindentation .............................................................................................................. 65 8.3 Scratch testing................................................................................................................... 76 8.4 Macroscopic wear testing ............................................................................................. 102

9 Discussion .................................................................................................................... 111

9.1 Mechanical properties of hardmetals......................................................................... 111 9.2 The scratch behaviour of WC-Co hardmetals........................................................... 114

Table of Contents ii

9.3 Abrasive wear ................................................................................................................. 126 9.4 Sliding wear .................................................................................................................... 127 9.5 Wear mechanisms: from the nanoscale and macroscopic scale............................. 128

10 Conclusions .............................................................................................................. 131

11 References ................................................................................................................. 133

1 Introduction 1

1 Introduction

Cemented tungsten carbide offers excellent wear resistance due to the combination of

the hard WC particles in a soft binder matrix making tungsten carbide one of the

oldest and most successful powder metallurgy products [1]. These hardmetals are

therefore used in a wide range of applications where wear resistance is very

important, such as sand blast/spray nozzles, seals in slurry pumps and component

parts in the oil industry [2,3].

Nanoscratch testing and nanoindentation measurements allow an evaluation of the

wear and mechanical properties of materials on the local scale [4,5]. In hardmetals

and cermets it is possible to measure the individual properties of the binder phase

and hard carbide phase and to possibly develop models which describe the

macroscopic mechanical deformation on the basis of the microscopic properties. The

application of nanoindentation methods in studying the properties of different

phases in composite materials separately has been shown very successfully on

nickelbase-superalloys [6].

In hardmetals the wear properties are the most important mechanical properties.

Nanoindentation techniques allow also quantitative measurements of lateral forces

during indenting and scratching. The relations between hardness and modulus of

elasticity and of friction and wear can be investigated on the microstructural scale.

The scratch resistance is not directly related with the hardness or other mechanical

properties. However, a deeper understanding of the relation between hardness and

scratch resistance is of interest. [7]

The fracture mechanisms can be studied with in-situ deformation experiments in a

AFM or SEM. With scanning probe techniques the microstructure including their

magnetic properties and the nanotribological properties can be studied in detail and

could lead to a better understanding of the microstructure property relations.

2 Objectives 2

2 Objectives

The main objective of this study was to investigate the wear properties of tungsten

carbide hardmetals from the nano to the macro scale. These investigations would be

used to determine the correlation between the macroscopic and nanoscale wear

mechanisms in tungsten carbide hardmetals and to understand the microstructural

influences on the wear.

This objective was explored through two main experimental approaches:

• The nanoscale wear of WC-Co hardmetals was conducted by carrying out

scratch tests on the samples .

• The macroscopic wear was conducted by carrying out sliding wear tests using

a tribometer and by three body abrasive wear tests.

The microstructural properties of the hardmetals were varied and two main features

investigated:

• The influence of the WC grain size on the wear mechanisms at both nano and

macro scale.

• The influence of the binder content on the wear mechanisms.

In addition the influence of the mechanical properties of the hardmetals on the

sliding wear properties will be investigated.

3 WC-Co hardmetals 3

3 WC-Co hardmetals

3.1 Introduction

WC-Co hardmetals are made by cementing very hard monocarbide grains (WC) in a

binder matrix of tough cobalt metal (Co) by liquid phase sintering. The high

solubility of WC in cobalt at high temperatures and the excellent wetting of WC by

the liquid cobalt binder result in optimum densification during liquid phase

sintering, producing a structure with little porosity [8]. The resultant cemented

carbide has high strength, toughness and hardness.

3.2 Hard metal manufacture

The manufacturing process consists of five main steps:

• production of powders,

• milling,

• pressing,

• pre-sintering and

• sintering.

Several techniques are available for the manufacture of tungsten carbide powders.

These include a traditional method based on the production of tungsten powder via

the hydrogen reduction of tungsten oxide followed by carburisation. The direct

caburisation of tungsten oxides can also be conducted [9]. The spray conversion

process allows the WC-Co powder to be produced in situ [10].

Milling of the powders is carried out to produce a homogeneous dispersion of

tungsten carbide in cobalt. During this process the particle size of the tungsten

carbide is normally reduced and stresses are induced in the particles which facilitates

the sintering process [11-14]. The cobalt may also undergo a phase transformation

during milling from a predominantly cubic structure to a hexagonal close packed

structure. A lubricant is added to the blended powder, mainly to reduce the friction


between the powder mixture and the surfaces of the tools and also to minimize the

tendency to form cracks [15].

The milled powders are pressed into shape using rigid steel or carbide dies with

pressures of up to 150-990 MPa [15]. Components can be pressed directly into

specified shapes or they may be pressed into large blocks which will later be shaped.

Following the pressing process the composite is pre-sintered in hydrogen with the

temperature increasing from room temperature up to 800 °C. The hydrogen reduces

the amount of adsorbed oxygen and oxides on the surface of the particles. When

cooled the material is coherent enough to allow further shaping and it is also less

susceptible to damage in the compacted form.

Sintering is normally conducted in vacuum at temperatures between approximately

1350 °C and 1550 °C. The rounded shape of the tungsten carbides in the early stages

of sintering leads to a facetted morphology which results in the flat trigonal prism

shape of the WC grains [16].

3.3 The constituent phases

3.3.1 Cobalt

Cobalt is the most commonly used binder for WC because of its excellent carbide

wetting and adhesion properties. The capillary action of cobalt during sintering

allows the achievement of high densities [8].

Cobalt exists in two allotropic forms, the hexagonally packed form which is stable at

temperatures below 417 °C and the face centred cubic form which is stable up to a

temperature of 1495 °C, the melting point of cobalt [17,18]. However, a significant

amount of fcc cobalt is present in sintered WC-Co hard metals at room temperature.

The transformation between the two phases is martensitic in nature [19,20]. The


cobalt transformation is also affected by the amount of tungsten carbide dissolved in

the binder and the binder mean free path, which is the distance between two carbide

particles. A high tungsten and carbon concentration in the binder has been shown to

increase the martensitic transformation temperature from 417 °C to approximately

750 °C [10]. This prevents the formation of the brittle hcp phase at low temperatures.

Hardmetals with a finer microstructure, have been found to have a higher fcc/hcp

ratio than conventional grades, this is due to the higher solution of tungsten in the

binder phase [10].

3.3.2 Tungsten carbide

Tungsten combines with carbon to form two carbides, WC, which has a maximum

microhardness of 24 GPa and W2C with a microhardness of 30 GPa [21]. Pure WC

does not melt under standard atmospheric conditions, but decomposes into a liquid

phase and graphite above a temperature of approximately 2780 °C as seen in the W-C

phase diagram.

Figure 3-1: Pseudo-eutectic WC-Co phase diagram [22]


The major phase in cemented carbides is the monocarbide WC. This has a simple

hexagonal crystal structure with two atoms per unit cell and a c/a ratio of 0.976 [16].

The W atoms occupies the 0,0,0 position and the carbon atom is located in the 31 ,

32 ,

21 or in the

32 ,

31 ,

21 position resulting in a non-centrosymmetric crystal structure

[23]. The WC crystal structure is polar with two sets of three equivalent { 0110 } planes

which leads to the formation of triangle-shaped crystals.

As a result of its crystal structure many of the properties of individual tungsten

carbide grains, including hardness, are highly anisotropic. However, no anisotropy is

observed in the sintered materials due to the random orientation of the grains.

Figure 3-2: The crystal structure of hcp tungsten carbide [21]

The hardness of tungsten carbide grains at room temperature varies from a

minimum of 10 GPa (measured parallel to the c-axis on the { 0110 } planes) to a

maximum of 24 GPa (measured along the basal plane) [20]. Up until the mid 1960s

WC was considered to a perfectly brittle phase however a lot of evidence exists

showing the plastic deformation of WC grains. Slip band formation has been

observed close to hardness indentations in WC single crystals. Glide bands and

dislocation networks show that the WC grains in WC-Co hard metals are plastically

deformed during compression testing. The slip planes are { 0110 } (prism planes) with


< 0001>, < 0211 > and < 3211 > as the preferred slip directions [24,25]. It has been

suggested that the { 0110 } planes are the most energetically favourable for cleavage.

3.3.3 The microstructural features of WC-Co alloys

The grain size of the WC powder used, the amount of cobalt and the processing

parameters such as the sintering time and temperature determine the microstructure

of the cemented carbide. The four features that characterise the microstructure are

the WC grain size, cobalt content, binder mean free path and contiguity.

Figure 3-3: A schematic representation of the WC-Co alloy illustrating the microstructural

parameters, where λCo is the mean free path of the binder phase, dWC is the WC grain size and CWC is

contiguity [26]

The WC grain size is defined as the mean linear intercept of the WC phase and this

can vary from ultra fine (200 nm) up to coarse grained (5 µm).

The mean free path is defined as the average thickness of the binder between the WC

grains and is dependent on the cobalt content and the size of the WC grains.

Contiguity is a measure of the continuity of the carbide skeleton existing within a

WC-Co alloy and is defined as the fraction of the total WC grain boundary surface

area that is taken up by the WC/WC interface. Contiguity decreases with increasing

binder content, decreasing WC grain size and is also dependent on the processing

history of the carbide.


The grain size, mean free path and contiguity all influence the mechanical properties

of WC-Co hardmetals and this is discussed in the following section.

3.4 Mechanical properties of hard metals

3.4.1 Hardness

The hardness of WC-Co hardmetals is affected not only by the hardmetal

composition but also by the level of porosity and the microstructure. The hardness of

tungsten carbide based hard metals has been extensively characterised and been

found to increase with decreasing cobalt content and decreasing WC grain size [27].

The highly constraining WC grains increase the yield strength of the cobalt binder

and make the hardness of the hard metals highly dependent on the binder mean free

path and the overall cobalt content.

Lee and Gurland expressed this dependence in terms of a Hall-Petch type relation

given in equation 3.1. Results from Roebuck et al. suggest that this dependence

becomes invalid when the WC grain size is below 0.3 µm and the hardness is

significantly lower than that predicted by Lee and Gurland’s model [28].

HWC = a + bd-1/2 (kg/mm2) E 3-1

The increase in hardness with decreasing binder mean free path has been shown to

be much more pronounced in ultrafine hard metals than in the coarser materials. An

exponential relationship between the hardness and mean free path of cemented

carbides was suggested by Gurland and Bardzil [29].


Figure 3-4: Dependence of the hardness of WC-Co hardmetals on the binder mean free path [29]

The hardness in WC grains is strongly anisotropic and Knoop hardness values

ranging from 2000 HK on the (0001) plane and 1050 HK for indentation on the {1010}

and {1011} planes have been reported [21,30,31].

3.4.2 Fracture toughness

Fracture toughness indicates the resistance of a material to fracture in the presence of

a sharp crack. The fracture toughness is measured by the critical stress intensity

factor, KIC. The KIC values are calculated from equation 3-2 [32].

KIC = A ∑

iia

HF E 3-2

Where H is the hardness of the material, F the indentation load and ai the length of

each Palmqvist crack. When all the quantities are expressed in SI units the constant A

is equal to 0.2784.


A plot of the toughness versus hardness exhibits the expected relationship, i.e. the

toughness decreases with increasing hardness. However for nanostructured hard

metals there is no decrease of the toughness with increasing hardness which implies

that different toughening mechanisms are present in conventional and

nanostructured materials. The fracture toughness and hardness of hardmetals vary in

different ways when the composition and/or microstructure of the material are

varied. There is some disagreement in the literature as to what type of relationship

between these properties exists [33-35]

Figure 3-5: Fracture toughness plotted against hardness for a range of WC-Co hardmetals with

differing grain sizes and binder contents [36]

Roeback and Almond found that as the cobalt volume fraction and WC grain size are

decreased the fracture toughness tends towards a limiting value of about 7 MPa

m1/2 [9], later work on very fine grained hard metals confirmed this trend [36,37].

These results suggest that further refinement in the WC grain size would lead to an

increase in hardness without any sacrifice in the fracture toughness of the hard

metal.


3.4.3 Young’s modulus

Young’s modulus for WC-Co hard metals usually lies in the range of 400 to 700 GPa

[38]. Okamoto et al. studied the relationship between Young’s modulus and several

microstructural parameters in WC-Co hard metals [39].

Figure 3-6: Relationship between Young’s modulus and Co content for samples with a WC grain size

of 20 µm [39]

Figure 3-7: Relationship between Young’s modulus and grain size for 10 wt% Co samples [39]


The Young’s modulus had a value of approximately 577 GPa for samples with a WC

grain size ranging from 3 to 20 µm and 523 GPa for samples with a WC grain size of

30 µm. Therefore no clear correlation between the WC grain size and Young’s

modulus could be established. However, the Young’s modulus was found to be

inversely proportional to the Co content as shown in figure 3-6. This suggests that

the Young’s modulus of WC-Co hard metals depends on the ratio of WC to Co.

3.4.4 The strength of WC-Co hardmetals

Cemented carbides have high compressive strength, much greater than those of most

other materials. Typical values of compressive strength range from 3.5 to 7.0 GPa

[38]. These materials have low ductility at room temperature so that the little

difference exists between their yield strength and fracture strength. At higher

temperatures the ductility increases slightly. The yield stress decreases

monotonically with increasing temperature and the fine-grained carbides tend to lose

their yield strength much more rapidly than the coarser grades. However at room

temperature the fine-grained carbides exhibit a high yield strength.

Figure 3-8: The effect of WC grain size on hardness and compressive strength of WC-Co hardmetals

[40]


3.4.4.1 Effect of the WC grain size

The transverse rupture strength (TRS) is a three-point flexure test used to measure

the strength of sintered materials and is commonly used in the testing of WC-Co

hardmetals. Work by Exner showed that the WC grain size has an influence on the

TRS of WC-Co hardmetals. In his work a maximum TRS at a WC grain size of 3 µm

was reported [40].

Figure 3-9: Influence of WC grain size on the transverse rupture strength of WC-6Co hardmetals [40]

3.4.4.2 Binder mean free path

The binder mean free path is dependent on both the cobalt content and the WC

particle size and can be used to describe the distance moved by a dislocation when

the binder is free of precipitates. The transverse rupture strength exhibits a

maximum when plotted against the binder mean free path. The maximum is a

transition point from predominantly brittle to predominantly ductile failure. When

the binder mean free path is low the strength is controlled by the fracture toughness

and therefore increases with increasing binder mean free path. When the binder

mean free path reaches higher values the hardmetal yields before failure and the

strength is mainly controlled by the yield strength of the hardmetal, which decreases

with increasing binder mean free path.


Figure 3-10: The influence of the binder mean free path on the strength of WC-Co cemented carbides

[40]

3.4.4.3 Effect of the cobalt content

Cemented carbides used in technical applications normally contain between 5 to 25

wt% cobalt. The hardness of the material decreases with increasing cobalt content

while the compressive strength reaches a maximum at 5 wt% Co and then drops

sharply when the cobalt content is increased [41,42]. On the other hand, the

transverse rupture strength improves with increasing cobalt content up to a

maximum at a cobalt content of approximately 20 wt%. The composition that

provides maximum strength depends on other variables such as the WC grain size

[41-43]. In some cases a maximum strength has been observed with a cobalt content

greater than 30 wt%.


Figure 3-11: Effect of Co content on the strength of WC-Co hard metals [40]

3.4.4.4 Contiguity

It is difficult to determine the effect of contiguity on the mechanical properties due to

other microstructural characteristics that must be taken into consideration. However,

the limited work that has been carried out shows increasing hardness with increasing

contiguity [40,44].

3.4.5 Nanoindentation of WC-Co hard metals

Gee et al. carried out nanoindentation tests on WC-Co hard metals and cermets to

investigate the mechanical properties of the individual phases in the materials [45].

The following curve was obtained for the WC-Co hard metal that was studied.


Figure 3-12: Load-displacement curves for 1 mN indentations on a WC-Co specimen [45]

The hardness values that were measured for the binder phase in the hard metal

ranged from 20-40 GPa and the carbide phase had a hardness of between 150 and 170

GPa. These values are approximately an order of magnitude higher than expected.

The difference in the measured and expected values was attributed to two factors.

Firstly the indenter geometry used in the calculations differed from the actual

indenter used and secondly the indentation size effect might have also played a role

in the measurements. This would account for the discrepancy in the values of the

hardness obtained in this study compared to macroscopic hardness of the constituent

phases.

4 Friction and Wear 17

4 Friction and Wear

This work focuses on the wear of cemented carbides therefore an understanding of

the fundamentals of the friction and wear processes that take place is essential. The

main focus of the work is the wear. However the friction behaviour is also of interest

and will be discussed briefly. This will be followed by a discussion of the different

wear mechanisms that occur in tribological systems and chapter 5 will concentrate on

the wear of WC-Co hardmetals.

4.1 Friction

Friction is defined as the resistance encountered by one body in moving over

another. This definition covers two important types of motion, sliding and rolling. In

both these types of motion a tangential force is needed to move the upper body over

the stationary counterface. The ratio between the frictional force F and the normal

load W is known as the coefficient of friction and is usually denoted by the symbol µ:

µ = F/W E 4-1

The magnitude of the friction force is normally described by the value of the

coefficient of friction which can vary from 0.001 to greater than 10. The coefficient of

friction normally lies in the range of 0.1 to 1 for most common materials sliding in

air [46].

The frictional force needed to initiate sliding is usually greater than that necessary to

maintain it and therefore the coefficient of static friction is greater than the coefficient

of dynamic friction. Once sliding is established, µ is found to be nearly independent

of the sliding velocity over a wide range for many systems although at high sliding

speeds, of the order of tens or hundreds of metres per second, µd falls with increasing

velocity.


4.2 Introduction to Wear

Wear is defined as the progressive loss of material from the surface of a solid body

due to mechanical action, i.e. the contact and relative motion of a body against a

solid, liquid or gaseous counterbody [47]. Wear occurs in many different situations,

for example, in piston rings, gears and in human body joints such as the knee and

hip. In all these cases widely varying wear conditions exist. Friction and wear are

both characteristics of the engineering system, which is called the tribosystem;

represented in Figure 4-1: Schematic of the elements in a tribosystem

The tribosystem usually consists of four elements,

1. A solid body,

2. Counterbody,

3. Interfacial element,

4. Environment.

The nature of the various elements that make up the tribosystem will clearly have a

direct effect on the wear process. The counterbody may be a solid, a liquid, a gas or a

mixture of these. Interfacial elements include lubricants, adsorbed and oxidised

layers and solid particles.

Wear can be divided into the following categories:

• Solid particle erosion

• Slurry erosion

• Cavitation erosion

• Abrasive wear

• Sliding wear

• Adhesive wear

In this work the focus is on abrasive, sliding and adhesive wear and these will be

described in further detail in the sections that follow.


Figure 4-1: Schematic of the elements in a tribosystem [47]

The material’s intrinsic surface properties such as hardness, strength and ductility

are also important factors that affect the wear resistance of a component. In addition

to the material properties other factors such as the surface finish, load, speed,

corrosion, temperature and properties of the counterbody also play an important role

in the wear process.

The mechanism of wear is very complicated and the theoretical treatment of wear

usually simplifies the processes that take place into four categories. These are

abrasion, adhesion, erosion and sliding, which may act individually or in

combination.

4.2.1 Abrasive Wear

In abrasive wear, material is removed or displaced from a surface by hard particles,

or by hard protuberances on a counterface, forced against and sliding along the

surface. Abrasive wear can be sub-divided into two types: two-body and three body

abrasive wear. Two-body abrasive wear is caused by hard protuberances on the

counterface, while in three-body abrasive wear hard particles are free to roll and

slide between two surfaces. Two-body wear is normally more severe than three-body

abrasion and may be one degree of magnitude greater [48-50]. This is because loose


abrasive particles only abrade the surface 10 % of the time and spend 90 % of the

time rolling [51].

Figure 4-2: Illustration of the differences between (a) two-body abrasion and (b) three-body abrasion

Material loss in abrasive wear occurs generally through several processes namely,

microploughing, microcutting and microcracking. In microploughing, an abrasive

particle pushes material in its path to both sides of the wear groove. Volume loss due

to the single passage of an abrasive particle does not normally occur but the repeated

action of many abrasive particles leads to the eventual removal of material as the

result of low cyclic fatigue [51]. Microcutting leads to the removal of material from

the surface by the formation of chips, shavings and fragments. The surface is worn

and a groove is formed whose geometry depends on the size and shape of the

abrasive particle. Cutting causes the most severe rate of abrasive wear in ductile

materials [52].

Microcracking can occur when highly concentrated stresses are imposed by abrasive

particles especially in the case of brittle materials. As a result large wear debriss are

formed as a result of crack formation, the subsequent growth and interaction of the

cracks. Ductile materials normally exhibit microploughing and microcutting as wear

mechanisms, which mechanism plays a more dominant role depends on the angle of

attack of the abrasive particle.


Figure 4-3: Schematic showing the mechanisms of abrasive wear in materials [51]

In brittle materials microcracking is a dominant wear mechanism which leads to

wear volumes larger than the groove formed, as the result of material loss caused by

crack propagation.

Many materials that are used for their wear resistance normally contain a hard phase

distributed in a soft matrix. The wear response of such materials depends on the size

of the hard phase regions in comparison to the scale of the deformation caused by

individual abrasive particles. The deformation refers to the width of depth of the

indentation caused by each particle and when this is substantially greater than the

size of the hard particle the material behaves like a homogenous solid. A finely

distributed hard phase leads to an increase in the flow stress of the matrix and hence

an increase in the wear resistance.

4.2.2 Sliding Wear

Sliding wear occurs when two surfaces rub against each other. In most practical

applications the sliding surfaces are lubricated in which case the wear that takes


place is termed, lubricated sliding wear, however in many laboratory investigations

the surfaces slide in air without lubricant. This type of sliding wear is called dry

sliding wear. The amount of material removal during sliding wear is dependent on

the load and sliding distance. Wear is usually measured by either removing the

specimen at regular intervals and weighing or measuring it or by continuously

measuring its position with an electrical or mechanical transducer and determining

the wear from its dimensions. The friction force during a wear test is determined by

measuring the tangential force on the specimen of the torque on a rotating

counterface. A continuous measurement of the friction coefficient allows any changes

in the sliding behaviour to be observed, these changes normally indicate a change in

the surface nature or topography.

The sliding velocity also affects the wear that takes place since the sliding velocity

affects the rate of frictional energy dissipation. Wear also depends on the nominal

contact pressure and wear transitions are commonly induced by changes in contact

pressure. The linear dimensions of the specimen are also important and other factors

to also be considered are the testing temperature, and in the case of lubricated

systems the viscosity of the lubricant.

4.2.3 Adhesive Wear

Figure 4-4: Schematic of adhesive wear

Adhesion is the formation and breaking of interfacial adhesive bonds e.g. cold-

welded junctions. This can take place when surfaces slide against each other. Sliding

leads to high local pressure between contacting asperities, which results in plastic

deformation, adhesion and the consequent formation of junctions locally. Relative

sliding between the contacting surfaces leads to the rupture of these junctions and


subsequent material transfer from one surface to the other, in addition to the

production of debris and material loss. The presence of a lubricating or oxide film

reduces the tendency for adhesion to occur [52].

5 Wear of WC-Co hardmetals 24

5 Wear of WC-Co hardmetals

5.1 Deformation behaviour of hardmetals

There have been many attempts to model the deformation behaviour of WC-Co hard

metals and the results of these studies will be discussed here. The discussion will be

divided into three sections. The first two sections will focus on the deformation

mechanisms that occur in the WC grains and the cobalt binder. Finally the combined

deformation of the alloy will then be discussed.

5.1.1 Deformation and fracture mechanisms in WC

WC grains are able to plastically deform without the occurrence of brittle fracture

due to the presence of several slip planes. Early work on bulk samples suggested that

the slip planes are of the { 0110 } type and the slip directions are < 0001>, < 0211 > and

< 3211 > [21,30,53]. These slip systems have been observed in the transmission

electron microscope and the operation of these systems is thought to be sufficient to

provide the five independent systems necessary to produce the shape changes [54]

needed to maintain continuity in deformed polycrystalline WC-Co. Unit dislocations

with the Burgers vectors < 0001 >, 31 < 0211 > and

31 < 3211 > have all been observed

and said to be glissile on a number of different planes. Furthermore in some regions

of both deformed and undeformed samples dissociated partial dislocations were

observed on {1010} planes which are defined by the following reaction

31 < 3211 > →

61 < 3211 > +

61 < 3211 >

Later work claimed that only the { 0110 } < 3211 > system which produces slip

equivalent to { 0110 } < 0211 > and { 0110 } < 0001 > is active in the WC and there only

four slip systems would be available [55]. This would limit the deformation of

polycrystalline WC and lead to interfacial crack formation. The occurrence of


microcracking at the WC/WC interface supports the lack of sufficient slip systems.

The number of slip systems in WC is therefore not very clear [55,56].

Deformation of WC grains by glide of 61 < 3211 > partials is usually associated with

slip in intense shear bands and intersecting bands often result in the nucleation of

microcracks [56-58].

WC grains usually have a relatively high dislocation density in the as-sintered

condition and can therefore plastically deform through dislocation multiplication.

5.1.2 Co binder deformation characteristics

The binder in as-sintered WC-Co hard metals is mainly present as fcc, this is unstable

at room temperature and results in the presence of numerous stacking faults in the

undeformed state. The energy in the binder is lowered by a fcc to hcp martensitic

transformation which takes place during deformation and is accompanied by twin

formation [59-61]. This is facilitated by the movement of 61 <112> partial dislocations.

Since the binder is a solid solution of W and C dissolved in Co the martensitic

transformation is possibly retarded by the dissolved W and C which lower the

martensitic transformation temperature. The transformation is accompanied with a

change in shape and the rigid WC skeleton restricts the transformation.

Since the fcc lattice is unstable at room temperature the stacking fault energy

becomes very low and therefore dislocations and stacking faults are the prominent

feature in deformed samples.

Thin lamellae of hcp material surrounded by fcc binder have been observed by x-ray

diffraction of deformed WC-Co hard metals. Lamellae of hcp material form readily in

the binder upon deformation and the stacking faults present in the undeformed

material act as nucleation planes for the lamella formation. During deformation the

stacking faults increase in thickness and eventually coalesce to form the lamellae of

hcp material that are observed. A complete transformation from fcc to hcp has never


been observed in these alloys. Work by Sarin and Johannesson estimated that less

than 10% of the binder transforms before fracture of the material takes place [62].

Therefore four types of defects are introduced in the binder phase during plastic

deformation [63]:

i. dislocations

ii. stacking faults

iii. twins

iv. regions where the face-centred cubic lattice is transformed to stable hexagonal

lattice

5.1.3 Deformation of WC-Co cemented carbides

Deformation in the hardmetal begins with the ductile deformation of the carbide via

a glide mechanism. The complex structure of the carbide skeleton means that a small

amount of plastic deformation in a single grain can result in large deformation in

other parts of the skeleton [62]. The cobalt binder then undergoes a transformation

into a hcp lamella. This is caused by the glide of partial dislocation on one of the four

slip planes of the fcc binder. The hcp lamellae only have one slip plane so that

deformation becomes increasingly difficult as it progresses. This leads to crack

formation in the binder phase and the simultaneous break up of the WC grains. The

dominant mechanism in this process is the break up of the carbide skeleton and is

observed by the large number of cracks at WC/WC interface as opposed to the

WC/Co interface.

The fcp to hcp transformation during the deformation of WC-Co toughens the

composite however as the deformation progresses this ability decreases and the

binder can no longer impede cracks [60].

5.1.3.1 Crack propagation in WC-Co alloys

Four types of fracture paths can be distinguished in WC-Co hard metals:


C transgranular fracture through the carbide crystals

C/C along carbide grain boundaries

B transgranular through the binder phase (Co)

B/C along binder/carbide boundaries

Fracture paths B and B/C are formed by the nucleation and coalescence of voids in

the binder phase. The voids are formed by the hydrostatic stress that develops in the

ligament as the binder is stretched without debonding from the WC grains [64]. The

work of fracture along the binder/carbide interfaces is smaller than for transgranular

fracture through the binder.

Carbide fracture normally precedes binder fracture and determines the direction and

type of path the crack takes. However the main contribution to fracture energy comes

from path B. The crack normally initiates in the brittle WC phase and avoids the

ductile binder which forms ligaments attatched to the WC grains. As the crack

continues to open each ligament is stretched until it ruptures. At this point the crack

tip in the matrix has moved further ahead. The local geometry of the microstructure

determines whether the crack advances along a B or B/C type path.

The plastic deformation of the binder is constrained when the surrounding WC

skeleton is intact. When the carbide next to the binder region cracks, localised

deformation of the binder begins under plastic strain. As the ligament is stretched its

lateral contraction is impeded by its continuity with the carbide and pore-like crack

blunting and/or the formation of voids inside the ligament occurs.


Figure 5-1: Schematic of the crack tip region in WC-Co during deformation. The black region

represents the binder phase, and the grey regions represent the WC grains. The white region is the

area where WC grain fall-out has occurred [65]

Crack paths B and B/C both lead to the formation of ligaments and voids in the

binder phase. The main difference is that the plastic deformation when the crack runs

parallel to the carbide/binder interface is reduced. Path B/C is therefore more

energetically favourable and is observed as long as the angle between the crack and

the interface does not exceed a critical value, which has been found to be about 25°.

When this value is exceeded the crack is forced to travel through the centre of the

binder region [64].

5.2 Sliding wear of WC-Co hard metals

Sliding wear occurs when two solid surfaces slide against each other and is similar to

abrasive wear in that they both require relative motion between two surfaces. Sliding

wear is more of a surface phenomenon which makes it difficult to predict the friction

and wear behaviour from bulk properties. The mechanical properties of hexagonal

materials are very anisotropic and this further complicates the wear behaviour of

WC-Co hard metals. Sliding wear is not a stable mechanism, during the course of

wear the contact surface changes and heat is produced because of friction and this

affects the wear rate. Pirso et al. carried out sliding wear tests on WC-Co hardmetals

[66]. During the initial stage of sliding wear no material detachment was observed

for the WC-6 wt% Co and the WC grains became glossy and polished. Carbide grain


fall-out was first observed after some distance had been covered, i.e. more than 1 km

of sliding distance.

Figure 5-2: Worn surface of WC-20 wt.% Co (a) WC grain size of 1.3 µm and (b) WC grain size of

1.5 µm, after 8 km run at load of 180 N [66]

In tests against a silicon nitride ball, WC-Co hard metal discs show an increasing

wear resistance with increasing hardness. The wear mechanisms have been found to

occur on a smaller scale than the individual WC grains. During sliding wear the wear

debriss are not easily removed and can therefore accumulate on the surface during

testing and form a tribofilm or mechanically mixed layer. The tribofilm has different

properties from the original material and may therefore influence the subsequent

wear behaviour. Tribofilm formation has been observed by Engqvist et al. on the

cemented carbides after dry sliding wear [67]. It has been suggested that small WC

fragments mixed with Co binder increase the surface toughness of the cemented

carbide and therefore also affect the wear resistance [67].

Cemented carbides have a high wear resistance in sliding contact, especially at high

normal loads. The wear mechanisms are best explained when the individual phases

are considered separately. The fragmentation of WC into small wear debriss is also

important for the relief of stresses in the material. These two features and the high


fracture toughness of the composite accounts for its ability to endure high pressures

without entering into a high wear regime.

A friction coefficient between 0.3 and 0.5 has been reported for cemented carbides

and the effect of the microstructure on the friction coefficient is still not clear [68].

Binderless carbides exhibit a high resistance to microfracturing and have similar

wear rates to cemented carbides containing a metallic binder phase. The main wear

mechanisms observed have been the fragmentation of WC grains and the oxidation

of the fragments.

Larsen Basse carried out sliding wear tests on WC-Co hard metals using a Rockwell

B indenter as a slider [69]. The cobalt binder was found to accumulate on the surface

during testing, similar observations were also made by Almond et al. [70]. It was

suggested that the binder is initially squeezed out of the surface by compressive

stresses in front of and on the sides of the indenter. Cobalt extrusion is followed by

cracking of the WC grains. This is followed by the microfracture of WC grains

adjacent to the surface defects, which occurs due to load concentrations around the

defect and reduced resistance to fracture as a result of binder flow and extrusion. WC

fragments are removed and cobalt smears on the surface of the hard metal. The

cobalt on the surface is expected to act as a lubricant which could reduce wear.


Figure 5-3: The variation in the wear rate with hardness of WC-Co hardmetals sliding against silicon

nitride at 9.8 N applied load and 31.4 mms-1 sliding speed [71]

Jia and Fischer carried out sliding wear tests on WC-Co hard metals with WC grain

size ranging from 1.5 to 0.9 µm [71]. The tests were conducted using a pin-on-disk

tribometer without lubricant. The volume loss in all the samples increased linearly

with sliding distance. The wear rate was found to increase with decreasing material

hardness (figure 5-3). The samples with 6 wt% cobalt exhibited a rapid decrease in

the friction coefficient with increasing WC grain size and a similar but less

pronounced dependence was observed for the softer materials. All the samples

exhibited a similar friction coefficient with a value between 0.4 and 0.5, even though

they had different wear rates. The WC grains on the worn surface were very smooth

and polished with occasional slip or cleavage features. Limited preferential binder

removal was observed and the wear scars showed no evidence of plastic deformation

on the micrometer or larger scale. In this work it was found that smaller WC grain

sizes led to a lower wear resistance in spite of an increase in hardness. The wear was

found to increase with increasing cobalt content, with the wear rate of

nanostructured materials with equal cobalt content being only 60% that of the


conventional counterpart. Jia and Fischer concluded that the best way to increase

sliding wear resistance was to reduce the cobalt content and increase the WC grain

size.

5.3 Scratch testing of WC-Co hard metals

Engqvist et al carried out scratch tests on single crystal WC crystals using a Vickers

diamond tip in air [72]. The loads were relatively high in comparison to the loads

used in this work. In this work a maximum load of 500 mN was used for the testing

but in the work reported form Engqivst a maximum load of 2 N was applied. The tip

was oriented with a corner in the scratch direction. The scratch behaviour of the WC

crystals was found to change according to the crystallographic orientation. The

grooves normal to the prism direction exhibited the highest scratch resistance and

the lowest amount of material removal. Ridges were formed along the scratch and

these were only observed for this crystal orientation. Scratch parallel to the prism

direction resulted in lowest scratch hardness and hence the highest material removal.

The amount of material removal varied between the different surfaces. The material

removal from the basal plane was approximately 3.5 times higher than from the

surface normal to the prism direction.

Shearing and cleavage along the scratch was found in the scratches on the basal

surface. Scratching in the orthogonal prism direction led to ridge formation and

cracks along the grooves. Strong material fracture was observed close to the indenter.

Slip bands were also visible and crack direction changed from being relatively

parallel to the scratch direction to being more aligned to the direction of the slip

bands. Scratching parallel to the prism direction led to crystal slipping parallel to the

scratch. No ridge formation or cracks were formed. Wear debriss were also observed.


Figure 5-4: Images of grooves in a) perpendicular prism and (b) parallel prism direction after scratch

testing with a Vickers diamond tip in air with a scratch velocity of 20 µm/s and an applied load of 2 N

[72]

Cracks perpendicular to the scratch direction were formed at the bottom of the

scratches on all the surfaces. Angular rod fragments were formed during scratching

in the orthogonal prism direction, these debriss were similar in nature to those

formed in the surface parallel to the prism direction. Scratching on the basal plane

produced debriss with a different morphology. The debriss had a triangular shape

presumably with the basal surface as the triangle plane. The region in front of the

scratch tip deformed plastically with slip band formation.

Jia et al also conducted scratch tests on WC-Co hard metals using a modified Vickers

hardness tester [71]. Single and multiple scratch tests were carried out with loads

varying from 1 N to 10 N and a speed of 0.05 mms-1. The sample with a WC grain

size of 2.5 µm and containing 10 wt% Co showed piled-up ridges, consisting of

extruded WC grains and binder material, along the scratch. A number of the

displaced WC grains were cracked and some exhibited slip lines indicating plastic

deformation. Some of the WC grains in the wear track were also cracked or

deformed. As the load was increased the WC grain cracking became more severe and

more material was piled up on the edges of the scratch. Less binder was observed in


the track suggesting that binder extrusion took place. The sample containing 6 wt%

Co with a grain size of 0.8 µm had a similar appearance after scratching. However

less material displacement was observed and there were also fewer WC grains

fragmented in the scratch. Additionally less binder extrusion was observed in the

materials with a finer microstructure.

The nano-structured hard metals studied in the work by Jia et al exhibited a higher

scratch resistance. The scratches were found to be smaller and this is correlated to the

higher hardness of these materials. No micro cracking was observed at loads of 100 g

and 500 g in the nano-structured materials. Cracks perpendicular to the scratch

direction were found at a load of 1000 g. The cracks are approximately two orders

larger than the grain size.

Work by Gee et al produced similar observations to the work by Jia and Fischer [73].

Scratch testing of WC-Co hard metals at loads ranging from 140-300 mN resulted in

debris formation and cracking of the WC grains. In addition to slip line formation,

binder extrusion and deformation along the edges of the scratches. Figure 5-5 shows

the surface of a WC-Co hard metal after a single scratch test with a 25 µm radius

Vickers diamond indenter and a load of 300 mN. Figure 5-5 shows the surface after

two passes over the same path.

Figure 5-5: Scratches on coarse grained WC-Co hardmetal, 300 mN applied load with a 25 µm

diamond indenter and a sliding speed of 0.1 mm/s (a) single pass and (b) 2 passes [73]


The damage observed by carrying out multiple tests over the same path indicated

that repeated damage may have a stronger contribution to wear than the initial

contact. This work showed that the wear mechanisms that occur on the macroscopic

scale also take place in micro-scale contact.

5.4 Abrasive wear of WC-Co hard metals

As a result of the different properties of the WC and Co phases the abrasive wear of

cemented carbides is complicated. The abrasive wear of WC-co hard metals is

divided into categories that depend on the size and hardness of the abrasive

particles. Wear by grits that are 1.2 times harder than the composite belong to the

“hard abrasive region” whereas the “soft abrasive region” refers to wear by softer

grits, which exhibit a lower wear rate [74]. Within each category there are sub-

divisions defined by the size of the abrasive grooves in relation to the microstructure.

If the grooves are large in comparison to the microstructure the wear is called tough

and if the groove size is comparable to or smaller than the microstructure the wear is

called mild. The abrasive wear of WC-Co hard metals will be discussed according

this nomenclature. The relative wear ranking of hard metals is highly dependent on

the triboysystem and is further influenced by the size and nature of the abrasive

used, the abrading wheel speed, the applied load and the material properties.

5.4.1 Mild abrasive wear

In mild abrasive wear the abrasive particles are affected by the individual properties

of each phase in the hard metal and are not affected by the average bulk properties.

The wear resistance of a composite can be determined from the load distribution on

its phases and their individual wear resistances [75]. The optimal wear resistance is

attained when the phases are worn down in parallel and the minimal wear resistance

occurs when each phase is worn independently of the other. In cemented carbides

the fraction of the hard carbide phase is high which means that the thin layers of Co

binder between the carbide grains can be assumed to behave differently from bulk

cobalt. Different cemented carbides with hardnesses ranging from 950-1850 HV have


been found to have very similar wear resistances, very similar to the average of

monocrystalline WC.

Figure 5-6: Schematic illustration of the theoretical optimum load and minimum load distribution

modes [76].

5.4.2 Hard abrasives

Large hard grits act as cutting tools and lead to the formation of grooves that are

larger than individual WC grains. This takes place in applications where the load is

high, such as the grinding of cemented carbides. The wear rate increases with the

groove depth in correlation to the indentation hardness. In this type of wear WC

grains are removed from the surfaces as small fragments and in some cases fatigue

cracks below the worn surface have also been observed [74].

5.4.3 Soft abrasion

In the soft abrasion region the abrasive particles are not able to penetrate the hard

metal surface. The abrasive particles slide over the surface pushing the WC grains

back and forth and causing the extrusion of cobalt and fragmentation of the carbide

grains. Damaged material is then either removed by the abrasive grits or smeared


over the material surface. In this wear regime the wear resistance has been found to

be proportional to the mean free path of the cobalt binder and finer WC grains and a

lower cobalt content increase the wear resistance.

The abrasive wear behaviour of fine-grained hard metals differs from that shown by

coarser grained hard metals. The fine-grained microstructures offer higher resistance

to microcracking than the coarser cemented carbides with the same hardness. As WC

grains decrease in size their individual fracture toughness increases, i.e. the crack

resistance increases [77-80]. The wear resistance was found to be dependent on the

WC grain size for a range of hard metals investigated by Quigley et al [77]. However

no dependence was determined for hardness below 1000 HV. All the grades

exhibited a linear relationship between abrasion and hardness. The coarser grades

had better wear resistance when the hardness was between 1000 and 1600 HV and

the finer grades exhibited a better wear resistance above a hardness of 1600 HV.

Figure 5-7: The variation of resistance with hardness of WC-Co composites from absrasion by

diamond, Full symbols: nanocomposites and open squares: conventional cermets [81].

The ratio between the WC grain size of the hard metal and the abrasive particle size

has been found to influence the material removal behaviour of WC-Co hard metals

during grinding by Hegeman et al. [82]. Earlier work by Anand et al. on the erosion


of WC-Co hard metals also indicated a dependency of the wear rate on the ratio

between the erodent particle size and the WC grain size [83]. Jia et al. observed a

similar relationship during scratch testing [71]. Therefore the wear mechanisms

observed in WC-Co hardmetals are not only dependent on the material

microstructure but also on the properties of the tribological system.

6 Tribological testing 39

6 Tribological testing

6.1 Sliding wear testing

Laboratory investigations of wear are carried out to either simulate practical

applications or to examine wear mechanisms, as in the case in this work. Many

different experimental arrangements have been used to study sliding wear. The

testing methods used to investigate sliding wear can generally be divided into either

symmetric or asymmetric arrangements. The most common asymmetric test rigs use

a pin pressed against a disc, either on a flat face or on the rim, a block loaded against

a ring or a pin on a flat. The load may vary from a few hundred millinewtons up to

several kilonewtons. Similarly the sliding speed can be varied.

Symmetric arrangements include the ring-on-ring tests with contact either along a

line or face to face. The most common asymmetric test set-ups use a ring pressed

against a disc, either on a flat face of or on the rim, a block loaded against a ring or

pin on a flat. The contact may initially be over an extended nominal contact area or

only at a point or line. In asymmetric arrangements the pin or block is normally

treated as the specimen and is the disc, flat or ring is called the counterface.

Figure 6-1: Geometries employed in sliding wear tests [52]


The friction can be measured continuously during a wear test by measuring the

tangential force on the specimen or the torque on a rotating counterface. This allows

changes in the sliding behaviour to be monitored and this is important since any

changes may indicate a change in the surface topography or a change in the wear

mechanism. The wear observed during sliding depends on the sliding distance and

to some extent on the sliding velocity and the duration of the test. The sliding

velocity affects the rate of dissipation of frictional energy and therefore affects the

temperature at the interface. The nominal contact pressure between the sliding

surfaces is important and another factor which may also play a role in the overall

wear process is the dimensions of the specimen.

6.2 Abrasive wear testing

Three body abrasion testing is of significant importance in the industrial application

of hardmetals. Several testing systems based on the ASTM G65 test have been

developed over the years for the testing of hardmetals [84]. The current experimental

setup that is discussed here was developed and built at the University of Erlangen-

Nürnberg and detailed information on the design and data capture system can be

found in the work by Herr [85].

The test setup is based on a combination of the ASTM G65 and ASTM B611-85 tests.

A specimen is pressed against a rotating steel wheel with a defined normal load that

is applied by a linear guide bar perpendicular to the specimen (figure 6-2). The wheel

is wet continuously with a liquid which can be a corrosive medium or water, as in

this work. The dry abrasive (SiO2, Al2O3 or SiC) is fed through a holding bin onto

the sample via a paddle wheel whose speed is controlled by a motor. As a result an

abrasive slurry is formed between the sample and the wheel so that the specimen

does not touch the wheel surface.


Figure 6-2: Schematic of the three-body abrasive wheel test apparatus designed at the University of

Erlangen-Nürnberg [85]

6.3 Scratch testing

During a scratch test the force normal to the sample is controlled and can be held

constant, increased or decreased at a linear rate. The scratch velocity and path

followed by the indenter are decided by the operator. The scratch velocity is usually

kept constant throughout the experiment and can range from 0.05 µm/s to 2.5 mm/s.

The tangential frictional force and the lateral scratch force are measured during a

scratch test, which allows the estimation of the scratch friction coefficient of scratches

made in any direction.


Figure 6-3: Schematic of scratch forces during a scratch test. The normal load is applied, the tangential

and lateral forces are measured and result from the material’s behaviour and indenter geometry [86].

A typical scratch experiment is performed in three stages: an original profile, a

scratch segment and a residual profile [86]. During the original profile stage, the

surface morphology is obtained by pre-profiling the surface under a very small load

at the location where the scratch will be performed. The actual penetration depth of

the indenter into the sample surface is determined by comparing the indenter

displacement normal to the surface during the scratching with the topography of the

original surface at each position along the scratch length. The roughness and slope of

the surface are taken into account in the calculation of the indenter penetration

during scratch segment. The residual scratch profile at one location is determined in

a similar manner. Additional morphological information can be determined by

profiling across the scratch, which can be incorporated into the overall test

procedure.

Under mild loading, scratches create elastic-plastic deformation which leads to a

groove with two adjacent lateral pile-up pads. These scratches are often described in

terms of the following parameters (Figure 6-4):

• Scratch width, a

• Scratch residual depth, p


• Scratch pile-up height, hb

• Pile-up height over scratch width, hb/a

• Scratch contact pressure, σ

• Scratch friction coefficient , fs

• Contact friction coefficient, fc

Figure 6-4: Cross profile of a scratch showing the main parameters used to describe a scratch [86]

Under severe abrasion conditions fracture processes can occur in which particles are

chipped out of the sample surface and/or cracks appear in and/or around the

scratch groove. This results in uneven scratch penetration and irregularities in the

tangential force curves and residual scratch morphology.

6.3.1 Importance of indenter geometry

The indenter geometry used in a scratch test influences the surface behaviour. There

are several indenter shapes available for scratch testing, the most common are:

• Conical indenter with a spherical tip

• Berkovich

• Cube corner

The tip shape used influences the deformation behaviour of the material and

additionally for pyramidal tips, the tip orientation during scratching also plays a


role. In the work by Youn et a.l the coefficient of friction was found to vary between

0.16 and 0.38 depending on the orientation of the Berkovich tip [87].

In this work a Berkovich indenter was used. It has a three-sided pyramid shape and

therefore the indenter orientation is important in testing. The pile-up formation is

less prominent when the edge of the Berkovich indenter is oriented in the scratch

direction.

6.4 Nanoindentation testing

Nanoindentation tests are generally conducted in order to determine the elastic

modulus and hardness of the specimen material from load-displacement

measurements. In nanoindentation testing the depth of penetration below the

specimen surface is measured as a load is applied to the indenter. The contact area of

the indenter can be calculated and this allows the modulus of the material to be

determined.

Figure 6-5: Schematic of the nanoindenter XP system used for nanoindentation and nanoscratch

testing [86]


6.4.1 Indenter types

Nanoindentation hardness tests are usually made with either spherical or pyramidal

indenters. The Berkovich indenter has a face angle of 65.03° and is normally used for

nanoindentation testing. The typical tip radius of a Berkovich indenter is in the range

of 50-100 nm.

Figure 6-6: Typical indenter geometries [88]

The cube corner indenter is mainly used for the indentation of ultra-thin films where

plastic deformation should be kept to a small volume. This is because it has a much

sharper angle giving it a smaller radius of curvature than the berkovich tip.

6.4.2 The indentation process

The indenter is driven into the material surface causing both elastic and plastic

deformation to take place resulting in a hardness impression in the shape of the

indenter. As the indenter is withdrawn only the elastic portion of the displacement is

recovered. This allows one to distinguish between the elastic and plastic properties of

the material. A schematic of indentation-load versus displacement data during one

cycle of loading and unloading is shown in figure 6-7.


Figure 6-7: Typical load-displacement curve for nanoindentation

The most important features are the peak load, the maximum depth, the final or

residual depth after unloading and the slope of the upper portion of the unloading

curve. Hardness and modulus are calculated from the load-displacement data for

each indentation by the Oliver and Pharr method [89].

6.4.3 Determination of the plastically deformed zone

The plastically deformed volume is given by the following equation

32 π ( 3ac)3 and 3ac ≈ 8.8h E 6-1

Figure 6-8: Schematic showing the plastically deformed zone formed by nanoindentation [89]


The hardness and modulus values are calculated from the load-displacement data for

each indentation according to the Oliver and Pharr method. In this method the

unloading curve is fitted to a power law with the form

P = Pm

n

m hhhh

⎟⎟⎠

⎞⎜⎜⎝

⎛−−

0

0 E 6-2

Where Pm is the maximum load, hm is the maximum displacement and h0 and n are

fitted constants. The plastic depth under load hp is determined from the following

equation:

hp = hm - εSPm E 6-3

where S is the stiffness (S = dP/dh), and ε is a constant (0.75 for a Berkovich

indenter). This equation was derived by Sneddon for a punch indenter pressed into

an elastic material [91]. If the area function for the indenter tip ((Ap = f(hp)) is known

the hardness H can be calculated from the following equation:

H = p

m

AP E 6-4

and modulus, E can be calculated from:

E = )( 22

112 ν

νπ

−⎟⎟

⎠

⎞

⎜⎜

⎝

⎛ −−

i

ip

EA

S E 6-5

Where Ei and νI are the modulus and Poisson’s ratio for the indenter and ν is the

Poisson’s ratio for the specimen being tested.

6.4.4 Indentation fracture mechanics

The contact loading of brittle solids can lead to elastic and plastic deformation in

addition to microcracking at and below the stressed surfaces. Cone-shaped Hertzian

cracks are an example of this kind of cracking. A circular cone-shaped crack

originates around the contact area between the sphere and flat surface of a brittle


solid when the critical load is exceeded. The crack grows from the circumference of

the contact area into the solid with increasing load. The maximum tensile stress

occurs at the contact area.

It is important to distinguish blunt and sharp indenters from each other when

looking at indentation problems. Pyramids or cones are considered as sharp and

spheres are considered to be blunt indenters. The type of indenter that is used

determines whether the contact is predominantly elastic or plastic.

Surface load by a point indenter results in median and lateral cracks below the

stressed surface [47].

Figure 6-9: Formation of median and lateral cracks in brittle solids due to indentation by a sharp

indenter [47]

An increasing point load leads to an increase in the size of the plastic zone around

and below the indentation. A median crack is formed when the load exceeds a critical

value. The crack grows in depth with increasing load and during unloading the


median crack is closed and lateral cracks are formed and propagate to the surface

under an applied load of less than F5. Reloading closes the lateral cracks and reopens

the median crack. Residual stresses due to plastic deformation cause the formation of

lateral cracks during unloading, these can play a part in the formation of microcracks

in the material. Residual tensile stresses lead to an increase in the crack length and

reduce the critical load required for microcracking. Median cracks may propagate in

depth due to residual stresses during unloading.

6.5 Microstructure analysis

The materials response to indentation or wear testing is determined by measuring

the amount of wear that occurs, or the indentation depth but in addition the

deformation of the microstructure is examined using electron microscopy. The

following sections describe how a scanning electron microscope works and also

describes a focused ion beam which can be used for further microstructural

investigation and advanced sample preparation.

6.5.1 Scanning electron microscopy

A scanning electron microscope (SEM) scans the sample with a high energy beam of

electrons in a raster pattern, these electrons interact with the sample surface

producing signals containing information about the sample surface topography,

composition and even the electrical conductivity.

A typical SEM normally has several detectors that are able to detect different signals,

which include secondary electrons, back scattered electrons (BSE) and characteristic

x-rays. Very high resolution images of the sample surface, with magnifications of up

to 100 000, can be obtained when using the standard detection mode, secondary

electron imaging (SEI). The high depth of field produced by the SEM result in an

image with a three-dimensional appearance.

In this work the initial electron microscopy work was conducted on a Hitachi S4800

cold field emission SEM (FE-SEM). A field-emission cathode in the electron gun of a

scanning electron microscope provides narrower probing beams at low as well as


high electron energy, leading to improved spatial resolution and minimised sample

charging and damage. The majority of the microscopic work was carried out on a

dual beam focused ion beam. The details of this are discussed in the following

section.

6.5.2 Focused Ion Beam

Figure 6-10: Schematic showing the basic operation of a focused ion beam [90]

A focused ion beam (FIB) is similar to a SEM however instead of using an electron

beam to image the sample it utilises a focused beam of gallium ions. Gallium wets

the tungsten filament and a big electric field results in ionisation of the gallium

atoms. It is possible to have a system that has both an electron and ion beam column.

This is called a dual beam FIB and with such a system the sample can be investigated

with either of the beams. A Zeiss 1540 EsB dual Focused Ion Beam (FIB) was used in

this work.

The accelerated Ga+ ions are destructive and sputter atoms from the material surface

and this allows the FIB to be used as a micro-machining tool and the sample surface

can be modified. Small sections of the material surface can be milled so that cross

section examinations can be carried out without having to section the whole sample.

The analysis can therefore be localised and the milled section can be immediately

observed using the electron beam in the case of a dual beam FIB:


6.5.3 Atomic force microscopy

An atomic force microscope (AFM) or scanning force microscope (SFM) is a very

high resolution scanning probe microscope, that produces a three dimensional

topographical image of the sample surface with a nano-scale resolution. An AFM

results in both excellent lateral and vertical resolution without complicated sample

preparation.

The AFM consists of a cantilever with a sharp tip at its end which is used to scan the

sample. The cantilever is usually silicon or silicon nitride with a tip radius of

curvature in the order of nanometers. When the sample and tip come into proximity

the cantilever is deflected due to the forces between the sample and tip. The

deflection is typically measured using a laser spot reflected on top of the surface of

the cantilever into an array of photodiodes. In most cases a feedback mechanism is

installed to avoid the risk of a collision of the tip with the surface that could occur if

the tip was scanned at a constant height. The feedback mechanism allows the tip-to-

sample distance to be adjusted to maintaining a constant force between the tip and

sample.


Figure 6-11: Schematic of an atomic force microscope

The AFM can be operated in several modes divided into two categories, static

(contact), and dynamic (non-contact). In the non-contact mode the cantilever is

vibrated. Most imaging is conducted using the contact mode, in this mode the force

between the sample surface and the tip is kept constant by maintaining a constant

deflection.

In the non-contact mode the cantilever does not contact the sample surface but

oscillates above the adsorbed fluid layer on the surface. Attractive van der Waals

forces decrease the resonant frequency of the cantilever, and are responsible for the

image formation [91].

7 Experimental methods 53

7 Experimental methods

7.1 Materials and sample preparation

The materials investigated in this study included a range of WC-Co hardmetal

samples, a binderless WC sample and pure cobalt. The binder content in the

hardmetals ranged from 6 to 15 wt% and the WC grain size ranged from 250 nm up

to 2.65 µm. The grades are classified as ultra fine, medium and coarse grained

depending on the size of the WC grains. The microstructure of the hard metals and

binderless WC sample is shown in the micrographs in figure 8-1. The nominal

compositions of the samples is given in table 7-1.

Table 7-1: Nominal composition and properties of the investigated WC-Co cemented carbides

Sample Binder wt%

Average WC grain size (µm)

Vickers hardness Hv

Fracture toughness MPa m1/2

CG15 15 2.65 1017 - UFG15 15 0.25 1486 - WC6M 6 1.21 1299 13.01

WC6MF 6.5 0.60 1449 10.25 WC6F 6 0.66 1591 9.30

WC6MG 6 0.48 1630 9.51 WC6SMG 6 0.25 1876 9.07

Scratch tests and nanoindentation measurements were carried out on polished

samples. A small section of the as-received sample was cut using a low speed saw

and the section was then polished. The specimens were polished using a Struers

RotoForce-4 automatic polishing machine combined with a Struers RotoPol-31 base.


The polishing discs and the polishing sequence used for the hardmetals are detailed

in table 7-2. The samples were cleaned in ethanol in an ultrasonic bath after each

polishing step.

Table 7-2: Polishing Sequence for the hardmetals

Step Disc Diamond Paste (µm)

Force (N) Time (minutes)

1 Struers MD Piano 120 H2O 15 5 2 Struers MD Piano 220 H2O 15 5 3 Struers MD Allegro 6 10 5 4 Struers MD Dac 3 10 10 5 Struers MD Dac 1 10 10 6 Struers MD Nap 0,4 10 30

7.1.1 Surface morphology of the materials

The surface morphology of the hard metal samples was examined using both an

atomic force microscope (AFM) and scanning electron microscope (SEM).

A Dimension 3100 Atomic Force Microscope (AFM) from Veeco, in contact mode,

was used to analyse the scratched samples and the section analysis function was

utilised to examine the scratch profile allowing for the scratch depth and width to be

determined. The tip frequency and velocity were kept below 1 Hz and 40 µm/s

respectively.

Initial scanning electron microscopy was conducted using a Hitachi S4800 FE-SEM.

An accelerating voltage of 10 kV or 15 kV was used and the working distance was

varied between 15 and 6 mm. The WC grain size was determined using the line

intercept method.

A dual beam Zeiss 1540 FIB was used to carry out further electron microscopy work

and FIB section analysis of the worn samples. Scanning electron micrographs were

obtained using an accelerating voltage ranging from 2 to 5 kV and the working

distance was varied from 1.5 to 5 mm.


FIB cross section analysis was carried out by milling a section across the length of the

scratch while the sample was tilted at an angle of 54°. A platinum layer was

deposited on the surface to protect it from damage that normally takes place when

the Ga+ ions hit the surface. The initial rough milling was carried out with a milling

current of 2 nA and conducted in two steps to avoid the curtaining effect that is

sometimes observed in milled surfaces, where a wavy morphology is obtained. The

fine polishing was conducted with a milling current of 200 pA to give a smoother

and finer surface finish.

7.2 Hardness and fracture toughness

The hardness of the materials was measured according to DIN ISO 3878 using a Leco

V-100A macrohardness tester with a Vickers hardness indenter. The diagonals

produced by the indenter were used to calculate the Vickers hardness. The

indentation diagonal measurements were carried out using an optical light

microscope. One specimen was tested for each material and an average of four

indents were made on each sample with a 30 kg load.

Figure 7-1: Micrograph of a Vickers indent on WC-6Co hardmetal showing the Palmqvist cracks at

the tips of the indent.


The fracture toughness of the materials investigated in this study was determined

from the lengths of the Palmqvist cracks using equation 3-2, which was initially

derived by Shetty et al. [32]. In figure 7-1 you can see a micrograph of a hardmetal

sample showing the crack formation.

KIC = A∑

iia

HF E 3-2

H is the hardness of the material, F the indentation load and ai the length of each

Palmqvist crack. When all the quantities are expressed in SI units, the constant A is

equal to 0.2784. The crack lengths were determined by measuring the end to end

distance between opposite cracks and subtracting the length of the diagonal of the

hardness indent. The measurements were carried out using a light microscope and

image analysis program, Image C.

7.3 Nanoindentation

Nanoindentation experiments were performed in the low and high load regime for

studying the local and global mechanical properties of the WC-Co hard metals

respectively. The hardness and Young’s modulus of the bulk materials and

constituent phase could be determined from the nanoindentation measurements.

The local properties were measured using a nanoindenting AFM (NI-AFM) from

Hysitron Inc. combined with a multimode AFM from Digital Instruments. This was

used to indent the WC grains and cobalt binder phases separately to measure the

mechanical properties of each phase in the hard metal. The indents were carried out

with a load of 5 mN. It was not possible to get a good resolution of the WC grains for

all the samples so limited results will be presented in the results section of this thesis.

The global properties were investigated at a higher load of 700 mN with a

Nanoindenter XP. Tip shape calibration was carried out using fused silica as the

calibration standard. A 4x4 array of indents, spaced 20 µm apart, were performed on


each sample with a Berkovich indenter using the continuous stiffness mode of the

Nanoindenter XP with a penetration depth of 2 µm. The Oliver/Pharr procedure is

incorporated into the MTS testworks software and was used to evaluate the

continuous stiffness measurements (CSM). With CSM measurements the hardness

and modulus of single indents is evaluated as a function of displacement.

7.4 Instrumented scratch testing

7.4.1 Nanoscratch testing

Nanoscratch tests on the hard metals were performed using a Nanoindenter XP with

a load controlled head. The load is applied normal to the sample surface via a

magnet/coil system which allows for precise and fast control. The indenter column is

supported by two leaf springs, providing very low stiffness to the vertical axis. A

maximum distance of 1.5 mm is allowed for the indenter travel normal to the sample

surface. Within this range the resolution is better than 0.1 nm. The maximum load

capacity for the standard system is 500 mN with a precision of less than 1 mN.

A diamond Berkovich indenter was used and the scratch tests were performed with

one corner of the Berkovich indenter in the scratch direction, this was important as

the tip orientation has been found to have an effect on the results [87,92].


Figure 7-2: AFM image of the Berkovich Tip used in the scratch testing

In this study the normal load ranged from 5 to 500 mN with a tip velocity of 0.1

µm/s. Tip calibration was performed using fused silica before each series of tests to

monitor the tip function and no significant change in the tip shape was observed.

The wear behaviour of the materials tested was characterised in terms of the scratch

width and depth. The scratch width and depth was measured using an AFM and

applying the cross-section function in the software. The wear mechanisms were

determined by examining the scratches using a SEM.

7.4.2 Microscratch testing

Microscratch testing or high load scratch tests were carried out on a MicroMaterials

Nanotest system in the Gordon laboratory of the University of Cambridge. The same

conditions that were used in the nanoscratch tests were applied to the microscratch

testing. Namely, a diamond Berkovich tip was used with one corner of the tip in the

scratch direction. The scratch velocity was maintained at 0.10 µm/s and the load

ranged from 1 to 10 N. Two single scratch tests and one multi-pass scratch test were


conducted at loads of 1 N, 5 N and 10 N on each sample and the scratch depth was

measured using a profilometer.

The nanoindenter system was located inside a sealed chamber as shown in figure

7-3 and this allowed the atmosphere during testing to be kept constant. The tests

were conducted under ambient conditions.

Figure 7-3: Photo of the nanoindenter set up in the Gordon laboratory, showing the vacuum chamber

in which the nanoindenter is enclosed [93]


(b)

Figure 7-4: A schematic of the nanoindenter from MicroMaterials [93]

7.5 Three body abrasive wear tests

Abrasive wear testing was conducted at room temperature on a three-body abrasive

wheel, which is described in detail in section 6.2. The wheel had a diameter of

195 mm, a thickness of 10 mm and was rotated at a speed of 100 rpm. The water flow

rate was 55 ml/min and the abrasive flow rate was 45 g/min. The abrasive used for

the experiments was SiO2 with a diameter of about 250 µm. A normal load of 100 N

was applied to the specimen.

The mass of the sample before and after testing was measured and the mass loss of

the sample was recorded. The mass loss was converted to a volume loss and the

volume loss was plotted against the wear distance for a wear test duration. The wear

distance (WD) is calculated using the wheel diameter (2r - mm), the wheel speed

(s-rpm) and the test duration(t - s), according to equation 7-1.


E 7-1

The sample geometry and resulting wear scar is shown in figure 7-5.

29mm 2mm 19.5mm

Figure 7-5 : Wear sample geometry: schematic [94]

Figure 7-6: Photograph of sample showing the as-received state on the left and the worn state on the

right [94].

7.6 Sliding wear tests

Sliding wear tests were conducted on a pin-on-disk tribometer, under dry conditions

at room temperature in laboratory air. The apparatus was built by the department of

mechanical Engineering at the University of Erlangen-Nürnberg. The WC-Co

hardmetals constituted the discs, which had a 40 mm diameter and were used in the

as-received condition without any additional polishing. The pin consisted of a steel

ball (100Cr6) with a diameter of ca 4 mm. The normal load applied was 9.8 N and the

6010002

⋅⋅⋅⋅

=tnrWD π


sliding speed was 0.10 m/s which translated to an angular speed of 477 rpm for a 10

mm track radius.

Figure 7-7 :Photograph of the (ball on disc) tribometer used for sliding wear tests

Two tests were completed for each condition. The wear coefficient k was measured

throughout the experiment. The wear profile was examined using a profilometer and

the samples were examined using a SEM.

8 Results 63

8 Results

In this chapter the results of the experimental work that was carried out are

described. The experimental data is presented in graphic form and electron

micrographs and AFM images are also displayed to show the microstructural

features of the wear processes that took place.

8.1 Material microstructure

The binderless WC had a multicrystalline microstructure with an average grain size

of 570 nm (figure 8-1a). The presence of pores was observed from the electron

micrographs of the sample. The pore shape and size varied with most pores

occurring on the grain boundary. The WC grain shape varied but the grains were

generally less angular than those in the WC-Co hardmetal.

The hardmetal samples were divided into two main categories, those with a low

binder content of 6 wt% and the second group containing the samples with a high

binder content of 15 wt%. The polished samples were examined using a scanning

electron microscope and the grain size was determined using the line intercept

method.

The WC grains were trapezoidal in shape with sharp edges and corners (figures 8-1b

to 8-1h). The WC grain size in each sample was found to vary and the size

distribution of the samples is shown in figure 8-3. These materials exhibited high

contiguity due to the low binder content. In some cases it was not possible to

distinguish the grain boundaries clearly. The average grain size of the samples varied

from 0.25 to 2.65 µm giving a wide range of grain sizes for investigation.

8 Results 64

Figure 8-1: Microstructure of the binderless WC sample and the seven WC-Co hardmetal samples

investigated in this work (a) binderless WC (b) WC6M (c) WC6MF (d) WC6F (e) WC6MG (f)

WC6SMG (g) CG15 and (h) UFG15

8 Results 65

0,01 0,1 10,0

0,2

0,4

0,6

0,8

1,0

Acc

umul

ativ

e Fr

eque

ncy

WC grain size (µm)

WC6MG WC6M WC6MF WC6F WC6SMG CG15 UFG15

Figure 8-2: Graph showing the grain size distribution in the WC-Co hardmetal samples

8.2 Nanoindentation

8.2.1 Nanoindentation of binderless WC

The mechanical properties of the binderless WC were investigated using

nanoindentation. The average hardness was found to be 27.8 GPa and the Young’s

modulus was found to be 736 GPa. Both the hardness and modulus were relatively

constant across the array of indents that were carried out as shown in the graphs in

figure 8-3.

0 10 20 30 400

5

10

15

20

25

30

35

40

Har

dnes

s (G

Pa)

Indent number

0 10 20 30 400

100

200

300

400

500

600

700

800

You

ng's

Mod

ulus

(GP

a)

Indent number

Figure 8-3: Hardness and modulus values for binderless WC

8 Results 66

Indentation of the binderless WC resulted in slip line formation in the WC grains as

can be seen in figure 8-4. In addition cracks were also observed in some of the WC

grains (figure 8-4b). Crack formation was predominantly transgranular however

crack formation along the WC/WC boundary was also observed.

(a) (b)

Figure 8-4: Micrographs of indents on binderless WC showing (a) slip line formation and (b)

intergranular grain fracture during nanoindentation

8.2.2 Nanoindentation of cobalt

0 2 4 6 8 10 12 14 160

1

2

3

4

5

Har

dnes

s (G

Pa)

Indent number

0 2 4 6 8 10 12 14 160

20406080

100120140160180200

Mod

ulus

(GP

a)

Indent number

Figure 8-5: Hardness and modulus values for the cobalt sample

The average hardness of the pure cobalt sample was found to be 3.30 GPa and the

Young’s modulus was found to be 132 GPa. Figure 8-5 shows the variation in the

8 Results 67

hardness and modulus values across the indent array and it can be seen that the

mechanical properties appear relatively constant.

8.2.3 Nanoindenation of WC-Co

The mechanical properties of the hardmetals were also determined by carrying out

nanoindenation experiments using a Nanoindenter XP. Figures 8-6 and 8-7 show the

typical load-displacement curves for the different hard metal samples. CG15

exhibited the highest indentation depth of about 1400 nm, whereas the lowest

indentation depth was exhibited by WC6SMG, with a value of about 1100 nm.

0 400 800 1200 1600 20000

100

200

300

400

500

600

700

800

Load

on

sam

ple

(mN

)

Displacement into surface (nm)

WC6M WC6MF WC6MG WC6F WC6SMG

Figure 8-6: Load displacement curves for WC-Co hardmetal samples containing 6 wt% binder

0 400 800 1200 1600 20000

100

200

300

400

500

600

700

Load

on

sam

ple

(mN

)

Displacement into surface (nm)

coarse-grained ultra-fine grained

Figure 8-7: Load displacement curves for samples containing 15% binder

8 Results 68

The variation in the hardness with WC grain size is shown in figure 8-8. The

hardness of the samples ranged from 12.7 GPa for CG15 to 25.5 GPa for WC6SMG. A

smaller grain size resulted in an increase in the measured hardness, which can be

correlated to the reduced indenter displacement. Secondly, a much higher hardness

was found for samples with a lower cobalt content.

0,0 0,5 1,0 1,5 2,0 2,5 3,00

5

10

15

20

25

30

35

Har

dnes

s (G

Pa)

WC grain size (µm)

6% binder 15% binder

WC hardness

Co hardness

Figure 8-8: Hardness of WC-Co hardmetal samples obtained from nanoindentation measurements

Young’s modulus values of the hardmetal specimens ranged from 550 to 750 GPa.

The highest value for Young’s modulus was displayed by WC6MF with a grain size

of 0.60 µm and a cobalt content of 6.5 wt%. A plot of Young’s modulus against the

WC grain size is shown in figure 8-9. From this diagram it can be seen that the WC

grain size does not have an influence on the modulus of WC-Co hardmetals.

However, Young’s modulus was found to increase with decreasing cobalt content,

i.e. Young’s modulus is inversely proportional to the Co content (figure 8-10). The

indenter size is comparable to the WC grain size and therefore the material response

is strongly influenced by the local mechanical properties of the material.

8 Results 69

The NI-AFM was used to perform indents on WC grains and cobalt binder separately

and thereby determine the hardness and Young’s modulus of the constituent phases

of the hard metal. Due to the limited resolution of the NI-AFM the mechanical

properties of the WC grains in three grades could be measured and the high

contiguity of the grades containing 6% binder meant that measurements in the

binder phase were limited to the CG15.

Figure 8-9: Influence of WC grain size on Young’s modulus

Figure 8-10: Influence of cobalt content on Young’s modulus

8 Results 70

Table 8-1: Indentation values for the WC phases in the different samples tested

Sample WC hardness (single

measurements) in GPa

Average hardness

23.59

25.16

WC6M

26.55

25.8±0.71

15.73

15.61

WC6MF

17.27

16.2±0.93

17.52

17.26

CG15

17.57

17.5±0.17

The hardness and Young’s modulus of the WC grains was found to vary in the

different samples tested. The WC grains in WC6M had an average hardness of 25.8

GPa compared with 17.5 GPa and 16.2 GPa in CG15 and WC6M respectively. In the

sintered alloy the carbide particles are randomly orientated so the bulk alloys does

not exhibit anisotropy, however individual tungsten carbide crystals are themselves

anisotropic. The anisotropic nature of the WC phases was further displayed in the

modulus values, with average values of 175, 256 and 375 GPa for WC6M, WC6MF

and CG15 respectively (see table 8-2). The average hardness of the cobalt binder in

CG15 was found to be 9.4 GPa and Young’s modulus was 342 GPa. The average

mean free path of the binder in this sample was 1.28 µm but this varied so that even

during indentation in the binder phase it was possible to encounter a carbide particle

as the tip moved into the surface.

8 Results 71

Table 8-2: Young’s modulus values of WC from hysitron measurements

Sample WC Young’s modulus

(single measurements) in

GPa

Average Young’s modulus

172.2

175.9

WC6M

177.6

175.2±2.76

261.9

238.7

WC6MF

267.4

256±15.23

321.7

307.6

CG15

318.1

315.8±7.32

The AFM micrographs in figures 8-11 show the indents that were performed on the

different hard metals with the same load of 700 mN, in figure 8.11d the individual

WC grains were indented with a load of 5 mN. The CG15 sample containing 15 wt%

cobalt exhibited cobalt lips (8-11a) on the edge of the indent due to the extrusion of

the Co binder. WC6SMG has a fine microstructure and exhibits pile-up behaviour

(figure 8-11c) which is a result of the deflection of WC grains which are pushed

upwards by the indenter.

8 Results 72

Figure 8-11: AFM images of indents on (a) CG15 (b) UFG15 (c) WC6MF and (d) WC6M, showing

the pile-up formation around the indents and the formation of cobalt lips in CG15.

The pile-up behaviour of the UFG grades is further illustrated in the topographical

image of UFG15 in figure 8-12. The pile-up formation is uniform on all three sides of

the indent indicating isotropic behaviour of the WC-Co hardmetal. In addition pile-

up behaviour is also observed in the indents formed on the individual WC grains

(figure 8-11d). In this case the pile-up formation was not always uniform indicating

the anisotropic behaviour of WC which is a result of its crystal structure.

8 Results 73

Figure 8-12: Pile-up formation around an indent in UFG15

Examination with a SEM showed that the WC grains at the bottom of the indent

contained glide lines indicating that slip was one of the mechanisms for plastic

deformation. Cracks were also observed in some of the WC grains, but the crack

formation was not extensive and no grains were completely fractured during

indentation. The microstructure within the indent appeared compacted as the WC

grains had been pushed closer together by the indenter and did not relax after the

indenter was removed. This pressing also led to the formation of cobalt lips by the

squeezing out of the binder.

8 Results 74

Figure 8-13: SEM Micrograph of an indent on WC6MF showing glide line and crack formation in the

WC grains.

It is also clear from figures 8-11a and b that the number of deformed grains during

indentation varies strongly from the coarse grained to the fine grained hardmetal.

The indent size in CG15 is much larger. However from figure 8-11a only a few grains

are deformed. In comparison, in the UFG sample (figure 8-11b) the indent size is

smaller but includes a large number of deformed WC grains. The number of

deformed grains is related to the plastically deformed volume which was calculated.

It was assumed that the WC grains were cuboidal in shape and therefore the number

of deformed grains was estimated by dividing the total deformed volume by the

volume of a single grain. CG15 had the largest deformed volume but the least

number of deformed grains because of the coarse grained structure. The UFG

materials had smaller plastically deformed volumes but due to the finer

microstructures the estimated number of deformed grains was significantly higher,

see figure 8-14. Therefore the deformation behaviour of the UFG hardmetals is a bulk

material behaviour throughout the entire indentation process whereas in the coarser

grained hardmetals the individual WC grains influence the deformation behaviour

when the tip first enters the hard metal surface.

8 Results 75

0,0 0,5 1,0 1,5 2,0 2,5 3,0

1000

10000

100000

Num

ber o

f def

orm

ed g

rain

s

WC grain size (µm)

Figure 8-14: Estimated number of deformed grains during nanoindentation

The mechanical properties of the hardmetal samples are summarised in table 8-3.

Table 8-3: Summary of the characterisation of all the samples

Sample Binder

wt%

Average

WC grain

size

(µm)

Description Vickers

hardness

Hv [30]

Indentation

hardness

(GPa)

Young’s

Modulus

(GPa)

CG15 15 2.65 CG 1017 12.7 555

UFG15 15 0.25 UFG 1486 19.7 561

WC6M 6 1.21 CG 1380 18.4 530

WC6MF 6.5 0.60 MG 1575 22.1 638

WC6MG 6 0.48 UFG 1760 21.8 476

WC6F 6 0.66 MG 1710 22.0 615

WC6SMG 6 0.25 UFG 1940 25.5 623

8 Results 76

8.3 Scratch testing

The results of the scratch testing experiments are presented and discussed in the

following section.

8.3.1 Friction behaviour

The friction force along the scratch was measured during scratch testing using the

lateral force probe allowing the scratch friction coefficient to be derived. Typical

curves for the variation in the friction coefficient with scratch distance are shown in

figure 8-15. After an initial run in distance the friction coefficient reaches a stable

value between 0.3 and 0.5. The fluctuation observed in the scratch friction coefficient

is due to the two phase microstructure and the subsequent surface roughness of the

material. The indenter moves across a hard and soft phase and the different

mechanical properties of the two phases lead to a variation in the friction coefficient

along the scratch. It was generally found that the fluctuation in the friction coefficient

is more prominent in the coarse grained samples than in the fine grained materials.

The small WC grain size in the fine-grained samples meant that the indenter tip

moved quickly between the two phases and the material therefore acted more

homogeneously. Furthermore, fine grained hardmetals were found to have a lower

friction coefficient of approximately 0.3 compared to 0.5 for the coarse grained

samples.

10 20 30 40 500,0

0,2

0,4

0,6

0,8

1,0

Scr

atch

fric

tion

coef

ficie

nt

Scratch distance (µm)

WC6M WC6SMG CG15 UFG15

Figure 8-15: Scratch friction coefficient for four of the samples tested

8 Results 77

The friction coefficient was not affected by the scratch load and remained constant at

all loads tested.

The scratch friction coefficient of the pure cobalt sample was found to be 0.3. This is

comparable to the value for the hardmetal and the binderless WC sample which had

a scratch friction coefficient of 0.4. Therefore no major differences were observed in

the friction behaviour of the different materials, however the analysis of the resultant

scratch depth and width showed significant differences in the wear behaviour of the

materials.

8.3.2 Binderless WC

Scratch tests were also carried out on binderless WC. This was polycrystalline with

an average grain size of 0.57 µm. The variation in scratch depth with increasing load

for the binderless WC sample is shown in figure 8-16. The scratch depths that were

measured for this sample were very small, due to the high hardness of the sample, a

scratch depth of 369 nm was measured for a single scratch test with a 500 mN load

and a multiple pass test at the same load resulted in a scratch depth of 1829 nm.

0 100 200 300 400 5000

250

500

750

1000

1250

1500

1750

2000

Scr

atch

dep

th (n

m)

Scratch load (mN)

Single scratch Multiple scratch

Figure 8-16: Variation in the scratch depth with applied load of binderless WC after single and

multiple scratch tests with a diamond Berkovich indenter and sliding speed of 0.10 µm/s,

8 Results 78

Scratch tests in the 1-10 N load regime resulted in much higher scratch depths. A

maximum scratch depth of 31.4 µm was measured for a multiple scratch test with a

10 N load. The maximum scratch depth at 1 and 5 N loads was 5.28 and 20.81 µm

respectively.

8.3.2.1 Wear mechanisms during scratch testing

Single scratches in the low load regime (5-10 mN) led to fine crack formation in the

WC grains as shown in figure 8-17. In addition some of the WC grains were chipped

and slip activity was also observed. However the damage was generally very limited

and the surface did not undergo a lot of deformation. Multiple scratch tests led to

further slip line formation and grooving of the WC grains caused by the repeated

movement of the indenter tip.

Figure 8-17: Micrographs of binderless WC showing grain fracture and slip line formation after (a)

single and (b) multiple scratch tests with a applied load of 10 mN. The scratch direction is from right

to left

Single scratch tests at medium load (100-500 mN) led to extensive crack formation,

leading to grain fracture and debris formation. Glide bands were also formed in the

grains, with slips lines being observed in both larger and smaller grains. Grain

removal was also observed (figure 8-18) and thin ligaments were observed in regions

where the material had not sintered properly, as can be seen in figure 8-18a. Multiple

scratch tests resulted in further grain fracture and removal. Wear debris was

compacted into regions in the microstructure where grains had been removed from

8 Results 79

the surface. No continuous mechanically mixed layer was observed across the worn

area.

Figure 8-18: Micrograph showing the damaged surface of binderless WC after (a) a single scratch and

(b) multiple scratch test with a load of 500 mN

Scratch tests in the high load regime (1-10 N) resulted in the predominantly brittle

fracture of the material. This led to the formation of very smooth fracture surfaces as

shown in figure 8-19b. The WC grains show a faceted surface with several smooth

and flat faces. The WC grains also underwent shear and glide bands can be seen in

some of the grains. The grain fracture was mostly transgranular, however

intergranular fracture was also observed. Multiple scratch tests resulted in very fine

wear debriss with a powder like appearance.

Figure 8-19: Worn surface of binderless WC after (a) a single scratch test with a load of 1 N and (b) a

multiple scratch test with a load of 5 N

8 Results 80

Due to its high hardness the scratch depth in the binderless WC is relatively low and

the dominant wear mechanisms involve brittle fracture of the WC grains. At low

loads the wear response was mainly ductile with very minimum crack formation and

no grain fracture taking place at all. With an increase in the load the mechanism

changed from a ductile to a brittle mechanism with extensive cracking and cutting of

grains. However it is important that the ductile behaviour of WC at high loads

should not be overlooked. The grains on the surface, or closest to the indenter

undergo brittle wear however those further from the contact zone underwent ductile

deformation in the form of slip line formation. The ductile behaviour was further

enhanced by the regions containing material that was not fully sintered, see figure 8-

20. This unsintered phase formed ligaments which acted somewhat like a binder in a

sintered hardmetal.

Figure 8-20: Micrograph of binderless WC after a single scratch test with a load of 500 mN showing

ligaments formed from unsintered material

8.3.3 Pure cobalt

The pure cobalt exhibited very high scratch depths due to its relatively low hardness

and high ductility. The increase in depth with increasing load is shown in figure 8-21.

The scratch tests were conducted at loads of 5, 10 and 100 mN.

8 Results 81

0 20 40 60 80 1000

250

500

750

1000

1250

1500

1750

2000 Single scratch Multiple scratch

Scr

atch

dep

th (n

m)

Load (mN)

Figure 8-21:Scratch depth for pure cobalt sample

A single scratch test at 5 mN led to a scratch depth of 29 nm and this increased to 50

nm for a multiple scratch at the same load. The scratch depth for a multiple scratch at

10 mN and 100 mN was 291 nm and 1683 nm respectively.

8.3.3.1 Wear mechanisms in pure cobalt

The cobalt sample exhibited ductile behaviour within the load range tested.

Figure 8-22: Scratch grooves formed on the surface of a cobalt sample after (a) single scratch and (b)

multiple scratch test with a load of 10 mN and velocity of 0.10 µm/s. The scratch direction is from

right to left

8 Results 82

Single scratches at low and high loads led to the formation of a smooth wear groove

with pile-up ridges along the edge of the scratch. Multiple scratch testing at both low

and high loads resulted in spallation along the scratch edges with the formation of

wear debris resulting from small fragments of the material being broken off.

It is expected that cobalt would exhibit ductile wear and this is what happened. Due

to its relative softness the cobalt sample exhibits relatively high scratch depths and

scratch testing at low loads results in the formation of wear grooves with material

pile up along the scratch edges. There is no crack formation observed and spallation

of the material is seen to take place after multiple scratching at low loads and after

single scratch tests at high loads.

8.3.4 Hardmetals with a low cobalt content

Figure 8-23 shows the increase in scratch depth with increasing load for the samples

containing 6 wt% cobalt binder. Single scratch tests resulted in a scratch depth

ranging from approximately 500 nm to 800 nm at a load of 500 mN. The lowest

scratch depth at this load was exhibited by the UFG sample WC6SMG and the

highest depth was shown by WC6M which has the largest WC grain size of 1.21 µm.

Multiple scratch tests, which consisted of twenty consecutive passes of the indenter

over the same position, resulted in a significant increase in the scratch depth. For

example, the scratch depth for WC6M increased from 795 nm for a single scratch to

1808 nm for a multiscratch test at a load of 500 mN. In comparison the sample with

the smallest grain size, WC6SMG (grain size 250 nm) had a scratch depth of 579 nm

and 1644 nm for single and multiscratch tests at 500 mN respectively.

8 Results 83

0 100 200 300 400 5000

100

200

300

400

500

600

700

800

900

Scr

atch

dep

th (n

m)

Load (mN)


Single scratch test

(a)

0 100 200 300 400 5000

250500750

100012501500175020002250

Scr

atch

dep

th (n

m)

Load (mN)


Multiple scratch test

(b)

Figure 8-23: Plot of scratch depth against applied load for the 6 wt% Co samples, (a) single scratch

tests and (b) multiple scratch tests

8 Results 84

The percentage increase in scratch depth from single to multiple scratch tests was

determined from the following equation.

% increase = 100×−

Single

Singlefinal

ddd

E 8-1

Where dfinal was the scratch depth after twenty passes and dSingle the scratch depth

after a single scratch test.

The percentage increase in scratch depth from single to multiple scratch test was

found to be different for each of the materials, for example an increase of 243% was

found for WC6M, with an increase in scratch depth from 298 nm to 1022 nm for

scratch testing at 100 mN. Compared to an increase of 178% for WC6SMG, with a

change in scratch depth from 227 nm to 632 nm at the same load.

In figure 8-24 the WC grain size is plotted against the scratch depth for both single

and multiple scratch tests at a load of 100 mN. For single scratch tests no clear

correlation between the WC grain size and scratch depth could be observed, however

a trend was observed in the multiple scratch tests. An increase in the WC grain size

leads to an increase in the scratch depth i.e. a decrease in scratch resistance.

However, WC6F did not follow this trend and showed a much lower scratch depth

than would be expected from the trend. This sample contains additions of VC and

Cr3C2 in the binder phase. These modifications to the binder are known to increase

the corrosion resistance of WC-Co hardmetals and could possibly also lead to an

improvement in the wear resistance [95].

8 Results 85

0,0 0,2 0,4 0,6 0,8 1,0 1,20

50

100

150

200

250

300

350

Scr

atch

dep

th (n

m)

WC grain size(nm)

Contains VC and Cr3C2 Single scratch test

(a)

0,0 0,2 0,4 0,6 0,8 1,0 1,20

250

500

750

1000

Scr

atch

dep

th (n

m)

WC grain size (µm)

Multiple scratch testcontains VC and Cr3C2

(b)

Figure 8-24: Plot of scratch depth against WC grain size at 100 mN for the 6 wt% Co samples (a)

single scratch tests and (b) multiple scratch tests

The results for the scratch tests carried out in the load range 1 to 10 N are presented

separately. This is because within this load range the scratch depth experienced by

the hardmetals was much higher than in the previous tests (5-500 mN) and the

scratch mechanisms will also be discussed separately. The scratch profile was

measured by the same indenter that was used to carry out the scratch test. The

displacement of the indenter relative to the original surface was measured. Figure 8-

25 shows the scratch profile for WC6M and WC6F after a multiple scratch test with a

8 Results 86

load of 5 N. This is a typical scratch profile that was produced during this high load

scratch testing. Due to the hard nature of the WC grains and the damage mechanisms

that took place the scratch depth varied extensively along the entire length of the

scratch resulting in a very jagged profile.

0 50 100 150 200 2500

5000

10000

15000

20000

Scra

tch

dept

h (n

m)


Pass 1 Pass 20

0 50 100 150 200 2500

4000

8000

12000

16000

Scra

tch

dept

h (n

m)


Pass 1Pass 20

(a) (b)

Figure 8-25: Scratch profile of (a) WC6M and (b) WC6F after scratch testing with 5 N load

WC6M WC6MF WC6F WC6MG WC6SMG0

4

8

12

16

20

Scr

atch

dep

th (µ

m)

1 N 5 N 10 N

Figure 8-26: Average scratch depth for single scratch tests between 1 and 10 N load

The average scratch depth for single scratch tests at 1, 5 and 10 N was plotted for the

WC-6Co samples in figure 8-26. The highest average scratch depth for a single

scratch test at 10 N was shown by WC6MG which had a scratch depth of

approximately 17.7 µm compared to 9.27 µm for WC6MF. No clear correlation

between the scratch depth and WC grain size could be determined as shown in

8 Results 87

figure 8-27 which plots the WC grain size against the scratch depth for single scratch

tests performed with a load of 5 N.

0,0 0,2 0,4 0,6 0,8 1,0 1,20

2

4

6

8

10

12S

crat

ch d

epth

(µm

)

WC grain size (µm)

Figure 8-27:Graph showing the variation in scratch with the WC grain size for single scratch tests

performed on the WC-6Co samples with a 5 N load

8.3.4.1 Wear mechanisms in WC-6Co hardmetals

The wear behaviour of the hardmetals in the low load regime (5-50 mN) is discussed

in this section. Single scratch tests between 5 and 10 mN led to scratch grooves being

formed across the sample as shown in the SEM images in figure 8-28. Small cracks

were also formed across the groove. The crack length was relatively small and did

not span the entire width of the groove. The cracks formed perpendicularly to the

scratch direction and were concentrated in the centre of the groove. Glide lines were

observed in WC6M (figure 8-28a). The grains showing slip were mainly within the

wear groove however some of the grains along the boundary of the scratch also

showed glide bands, indicating that the WC grains outside the scratch were also

plastically deformed during scratching.

8 Results 88

Figure 8-28: Micrographs showing the wear mechanisms on the WC-Co hardmetals after (a,c,d) single

and (b,d,f) multiple scratch test at a load of 10 mN. The images of WC6M show slip line and crack

formation. Grain fall-out can be seen in WC6MF and a tribofilm is on the surface of WC6SMG after

multiple scratch testing. The scratch direction is from right to left

Multiple scratch tests led to the deepening and widening of the scratch groove. In

single scratch tests there was only one groove channel visible however with multiple

scratch tests the wear scar consisted of several channels formed by the repeated

movement of the diamond indenter across the surface. The cracks were still relatively

8 Results 89

small in length, however debris formation indicated that crack growth and

intersection does took place. Furthermore, cracks were also visible not only towards

the centre of the wear groove but along the edges of the groove. Glide activity was

now visible in all the samples. In addition to grain fracture grain removal was also

observed in some of the samples. Two large holes can be seen in WC6MF, these holes

show the positions where WC grains were pulled out of the hardmetal surface. In

this sample the wear debris, consisting of small WC grain fragments, was also

observed to sit between the WC grains, occupying the spaces previously occupied by

WC grains. In the case of WC6SMG a tribolfim was observed on the surface after

multiple scratching. This consisted of cobalt binder that was squeezed out by the

indenter during scratching. The WC grain fall out in this sample was much less

pronounced in comparison to the other samples and furthermore there was less

obvious crack formation.

Single scratches at medium loads (100-500 mN) resulted in extensive WC grain

cracking. Crack formation was concentrated along the edges of the scratch groove

and resulted in grain removal on the edges of the scratch. Some of the cracks also

extended across the entire width of the scratch. Pile-up formation on the edges of the

scratch took place (figure 8-29). In addition, glide band formation in the grains within

and adjacent to the worn areas could be seen, both in the large and in the smaller WC

grains.

Figure 8-29: Pile up in WC6M during single scratch at 100 mN

8 Results 90

Prismatic slip was observed in some of the larger grains and was identified by the

slips lines formed in three directions forming a triangular pattern (figure 8-30). Crack

formation in the WC grains occurred both in the slip direction as well as normal to

the slip direction.

Figure 8-30: Prismatic slip in WC6M after a single scratch test with a 200 mN load

Examination at high magnification allowed the deformation of the cobalt binder to be

carefully examined. The cobalt binder was found to have undergone plastic

deformation via void formation resulting in thin ligaments of cobalt loosely attatched

to the WC grains (figure 8-31). Ultimately the binder attains a porous-like structure

during scratching which makes it easier for WC grains to fall out. In figures 8-31a

and 8-31b WC grains can bee seen which are attached to thin cobalt ligaments. In

figure 8-31b the cleavage of the WC grain along the slip plane can also be observed.

The fracture surface of the WC grain in this micrograph indicates both ductile and

brittle wear mechanisms. The round particles present on the fracture surface are

thought to be unmelted WC particles from the sintering process.

8 Results 91

Figure 8-31: micrographs showing the worn surface after single scratch tests performed with 500 mN

load on various WC-6Co hardmetal samples.

Multiple scratches at high load resulted in crack growth and intersection leading to

the formation of a mechanically mixed layer on the wear surface. This is often

referred to by some researchers as a tribofilm [67]. This consisted of WC fragments

and cobalt binder mixed together (figure 8-32). This layer was strongly adhered to

the hardmetal surface and was not removed when the specimen was cleaned in the

ultrasonic bath after testing. This mechanically mixed layer covered the entire length

of the wear groove so that it was not possible to examine the underlying

microstructure without carrying out cross-section analysis. Because the scratches

formed were very fine (µm range) cross section analysis could only be carried out

using a FIB. Therefore the cross section examination of the hardmetals was done on a

limited number of samples. Figure 8-33 shows a cross-section of a scratch on WC6MF

after a multiple scratch test with a load of 500 mN. Cracks from the surface extend

into the bulk material as indicated in the marked region on the micrograph. The

mechanically mixed layer on the surface of the scratch is also visible and furthermore

8 Results 92

pores in the material can also be observed. These could accelerate crack growth into

the bulk microstructure.

Figure 8-32: Mechanically mixed layer or tribofilm on WC6MG after a multiple scratch test with a

load of 100 mN

Figure 8-33: A FIB cross-section of a scratch on WC6MF formed after a multiple scratch test with a

load of 500 mN.

Single scratch tests in the 1-10 N load regime resulted in widespread WC grain fall

out along the edges of the scratch as seen in figure 8-34a which shows the lower edge

of scratch on WC6M after a single scratch with a 1 N load. There was also extensive

8 Results 93

grain fracture throughout the worn region indicating predominantly brittle wear.

However glide lines were also observed on the grains adjacent to the scratch

showing the ductile deformation of the area not immediately in the contact zone of

the indenter. There was a build up of material at the bottom of the scratch, forming a

ridge, as a result of material being pushed by the indenter during scratching (figure

8-34b). There was also pile-up of the cobalt binder along the edge of the scratch, this

was further aided by the removal of the WC grains which made it easier for the

binder to be extruded. Multiple scratches led to further grain removal and fracture,

with the formation of wear debris which became embedded into the hardmetal

surface (figure 8-34c). The debris formed during multiple scratch testing was much

finer than that formed by single scratches (see figure 8-34d).

Figure 8-34: (a) Grain fall-out on the edge of a scratch performed with a load of 1 N and (b) material

pile-up in the scratch direction after a multiple scratch test with a load of 5 N. The scratch direction is

from right to left.

8.3.4.2 Summary of scratch testing on WC-6Co hardmetals

During a scratch test the load is first applied and the indenter penetrates the sample.

Once the testing load is reached the indenter moves across the sample at the selected

8 Results 94

velocity. At low loads the penetration depth of the indenter is very low and it only

scratches the surface of the material which leads to cracks and grooves in the WC

grains. Multiple scratching at low loads leads to fracture of the WC grains as a result

of crack growth and intersection. The fractured grain segments may remain

embedded in the material or be removed from the surface. Grain fall-out occurs due

to the movement of the sharp indenter tip which repeatedly pushes against grains

which are loosely anchored as a result of binder extrusion. Therefore initial extrusion

or deformation of the binder is critical. Grain fall out in the UFG samples was very

limited. This is because these materials are harder and therefore the penetration

depth of the indenter is lower and the damage mainly occurs on the surface of the

material. In addition, the low binder mean free path means that the indenter is likely

to encounter a hard WC particle when it tries to penetrate the surface keeping the

penetration depth low.

As the applied load is increased the penetration depth of the indenter is increased

and the damage occurs deeper in the material surface. The increased load leads to

more brittle deformation of the WC grains as the WC grains are no longer able to

absorb the load. At loads between 5 and 10 mN very small cracks were formed

towards the middle of the scratch groove The cracks formed approximately

perpendicular to the scratch direction and at higher loads crack formation occurred

on the edges of the scratch groove. Crack growth was mainly transgranular and in all

directions. Cracks were observed both parallel to the glide plane or normal to the

glide plane. The cobalt binder also underwent plastic deformation, forming thin

ligaments and voids. The WC grains are no longer fully attached to the binder and

this makes WC grain removal much easier.

When the load range was changed from 500 mN to 1 N the scratch depth was

observed to increase significantly in all the hardmetals. However the scratch

mechanisms observed were similar to those seen in the lower load regime. There was

more pile-up formation observed in this load range, with the indenter pushing

material along as the scratch progressed. There was also much more extensive WC

grain removal which led to extrusion of the binder phase which facilitated the further

8 Results 95

loss of WC grains. This was especially evident for the single scratch tests where the

edge of the scratch mainly consisted of binder with no WC grains (figure 8-33a).

What was interesting is that the apparent influence of the grain size on the wear

behaviour could not be observed in the macro scratch tests (N range).

8.3.5 WC-Co hardmetals with a high cobalt content

The variation in scratch depth with increasing load for the 15 wt% Co samples is

shown in figure 8-35. A scratch depth of 1079 nm was measured for UFG15 for a

single scratch at a load of 500 mN. In comparison CG15 exhibited a lower scratch

depth for the same test, with a depth of 781 nm being measured. The UFG sample

displayed a consistently higher scratch depth at all tested loads.

0 100 200 300 400 5000

400

800

1200

1600

2000

2400

2800

Scr

atch

dep

th (n

m)

Load (mN)

UFG single scratch UFG multiple scratch CG single scratch CG multiple scratch

Figure 8-35: The variation in scratch depth with load for the 15 wt% Co samples

Multiple scratch tests caused a significant increase in scratch depth for both samples

with the UFG sample once again exhibiting a higher scratch depth for multiscratch

tests at all loads. For example the change in depth from a single to a multiple scratch

test at a load of 500 mN is 781 nm to 1.598 µm for the coarse grained material and

1.079 µm to 2.440 µm for the UFG sample. This is an increase of 126% for the UFG

material and 105% for the coarse grained sample.

8 Results 96

0,0 0,5 1,0 1,5 2,0 2,5 3,00

200

400

600

800

1000

1200

Scr

atch

dep

th (n

m)

WC grain size (µm)

Single scratch Multiple scratch

Figure 8-36: Variation in scratch depth with WC grain size for the WC-15Co samples

In this load range (5-500 mN) the scratch depth was found to decrease with

increasing WC grain size i.e. the scratch resistance was determined to be inversely

proportional to the WC grain size (figure 8-36). A finer grain size led to a higher

scratch depth.

In the 1-10 N load range the scratch depth increased drastically for both samples. The

scratch profile for single and multiple scratches on CG15 and UFG15 after scratch

testing with a 1 N load are shown in figure 8-37. The scratch depth in CG15 was

found to be relatively even however in UFG15 it varied along the length of the

scratch. The average scratch depth for single scratch tests is shown in figure 8-38.

UFG15 exhibited a lower scratch depth than CG15 in this load range. This is the

opposite of what was determined in the 5-500 mN load range where CG15 exhibited

better wear resistance and had consistently lower scratch depths than UFG15.

8 Results 97

0 50 100 150 200 2500

2000

4000

6000

8000

10000

12000

Scr

atch

dep

th (n

m)


Pass 1 Pass 20

CG15

0 50 100 150 200 2500

1000

2000

3000

4000

5000

6000

7000

8000

Scr

atch

dep

th (n

m)


Pass 1 Pass 20

UFG15

(a) (b)

Figure 8-37: Scratch profile after scratch testing at 1 N load. As measured with the scratch tip along

the scratch length (a) CG15 and (b) UFG15

0 2 4 6 8 100

4

8

12

16

20

Scr

atch

dep

th (µ

m)

Load (N)

CG15 UFG15

Figure 8-38: The average scratch depth after single scratch tests at loads between 1 and 10 N

8.3.5.1 Wear mechanisms during scratch testing

The micrographs in figures 8-39a and 8-39b show the material microstructure of

UFG15 after a single and multiple test with a 5 mN load. A single scratch resulted in

cracking of the WC grains. As a result of the very small grain size the cracks

extended across the entire grain leading to the fracture and chipping of the WC

grains. Grain fall out of the smaller WC grains was also observed after single

scratches at low loads (figure 8-39a). Multiple pass tests led to increased material

8 Results 98

removal and the formation of a mechanically mixed layer on the material surface

(figure 8-39b).

Grooving of the WC grains was observed in CG15 after single scratch tests at low

loads (figure 8-39c). Crack formation was also observed, the cracks were

concentrated in the centre of the scratch groove and formed normal to the scratch

direction. Cracks did not form in all the grains along the length of the scratch. The

pileup of the cobalt binder could also be observed and this led to the formation of

cobalt lips. Multiple scratch testing resulted in chipping of the WC grains due to

crack growth and intersection. This led to the formation of fine wear debris. The

chipped WC grain fragments were re-embedded into the binder phase and the pile-

up ridges consisted of cobalt binder containing WC grain fragments (figure 8-39d).

Crack formation during multiple scratch tests was still very limited.

Figure 8-39: Micrographs of UFG15 and CG15 after (a,c) a single scratch test and (b,d) multiple

scratch tests with a load of 5 mN, showing grain fracture and triboflim formation in UFG15 and

binder extrusion in CG15

8 Results 99

Scratch testing at medium loads (100-500 mN) led to more severe WC grain fracture

and removal in the UFG15, which correlates to the high scratch depths that were

measured. Cracks extended across the scratch width and grain removal along the

edges of the scratches was observed (figure 8-40). The marked area in figure 8-40a

shows a region on the edge of the scratch where WC grains have been removed after

a single scratch with a 500 mN load. Voids in the underlying binder could also be

observed. Multiple scratching in UFG15 resulted in further grain removal and the

formation of the mechanically mixed layer which was also observed at multiple

passes at low loads.

The typical wear damage observed in the CG15 after single scratch tests at medium

loads is shown in figure 8-40c. This image shows the formation of slip lines in the

larger WC grains next to the scratch. In this micrograph the alignment of the slip

lines in two adjacent grains is shown. Cracking, cutting and chipping of the WC

grains was also observed which led to the loss of material. In addition, binder

extrusion occurred. Furthermore, the pile-up of material along the scratch edges was

observed and is shown in figure 8-40d. This image shows a single scratch on CG15

performed with a load of 500 mN with ridges of material on the bottom edge of the

scratch. Multiple scratch tests resulted in further material loss as a result of grain

fracture and removal. A mechanically mixed layer was also observed on the worn

surface.

8 Results 100

Figure 8-40: Micrographs of UFG15 and CG15 after (a,c) a single scratch test and (b,d) multiple

scratch tests with a load of 500 mN. The marked area in UFG15 shows binder deformation and grain

fall-out.

The dominant wear mechanisms for single scratch tests in 1-10 N load regime were

WC grain fracture and removal. In CG15 widespread glide activity was also

observed with slips lines present in quite a number of the grains. This was less

pronounced in UFG15. In both cases multiple scratch tests led to increased grain

fracture and removal. The FIB cross sections in figure 8-41 show the mechanically

mixed layer formed during multiple scratch testing with a load of 10 N. A network of

cracks extending into the underlying material can also be seen and the intergranular

crack is observed.

8 Results 101

Figure 8-41: FIB cross section on a scratch formed after a multiple scratch test with a 10 N load. The

marked area is magnified in the right image, showing the extension of cracks into the subsurface of the

hardmetal (a) UFG15 and (b) CG15.

In CG15 the intergranular crack growth is more prominent and the mechanically

mixed layer is relatively thin in comparison to that formed on UFG15. In CG15 the

mechanically mixed layer is composed of grain fragments and in the UFG

counterpart whole grains are present in this layer.

8.3.5.2 Summary of scratch testing on WC-15Co hardmetals

At low scratch loads the coarse grained material exhibits very low scratch depths and

the damage is confined to grooving of the WC grains and extrusion of the cobalt

binder. The UFG material on the other hand shows much more severe damage

mechanisms. Single scratches resulted in grooving, cracking and chipping of the WC

grains and also the fall out of WC particles in the UFG sample. In CG15 the grains are

very large, increasing the probability that the indenter encounters a WC grain and

therefore only the surface of the WC grains is damaged. In comparison the WC grain

size in the UFG sample is comparable to the penetration depth at low loads and the

8 Results 102

movement of the indenter across the material can lead to the removal of the WC

grains. At a load of 500 mN a scratch depth of 781 nm was observed for the CG

material after a single scratch test which is very small in comparison to the average

grain size of 2.65 µm. On the other hand the scratch depth for the UFG sample at the

same load was 1.079 µm which is more than four times the average grain size of the

material.

At higher loads a similar mechanism would be expected i.e. a higher penetration

depth which allows the fine grains in the UFG material to be dug out by the indenter

and the larger CG grains are still much larger than the penetration depth and

therefore not easily removed. However at higher loads the force of the indenter

pushing against the WC grains also leads to the extrusion of the binder which

precedes the removal of WC grains from the surface. An increase in the load also

leads to crack formation in the grains which can no longer undergo plastic

deformation to absorb the load. The presence of slip lines and sharp fracture edges

was observed within single grains in CG15. Slip activity was not observed in UFG15,

though it is likely that it took place but due to the fine grain size was difficult to

observe.

8.4 Macroscopic wear testing

Sliding wear and three body abrasive tests were conducted on the WC-6Co

hardmetals to investigate the macroscopic wear behaviour of the materials. The

results of these tests will be discussed in the following section beginning with the

results of the three body abrasion testing.

8.4.1 Three body abrasive wear

Three body abrasive testing was carried out at 100 N load and with a rotation speed

of 100 rpm. The average size of the SiO2 particles was 200 µm. The results of the wear

loss for each sample is shown in figure 8-42.

8 Results 103

0

10

20

30

40

50

60

70

80

Vol

ume

loss

(10-3

m3 )

15 min 30 min 60 min 120 min

WC6M WC6MF WC6F WC6MG WC6SMG

Figure 8-42:Volume loss of WC-6Co samples after abrasive wear testing

The highest wear loss was shown by WC6M and the least wear loss by WC6SMG. It

was generally found that an increase in the WC grain size led to an increase in the

wear loss. There was an approximately six fold difference in the volume loss

observed in the UFG WC6SMG sample and coarse grained WC6M sample. The

change in the wear rate (defined as the volume loss per metre) with sliding distance

was also monitored and this is shown in figure 8-43. From this diagram it can be seen

that the wear rate is relatively constant for all tests on the same sample.

0,000

0,001

0,002

0,003

0,004

Wea

r rat

e (m

³/m)

15 min 30 min 60 min 120 min


Figure 8-43: Wear rate of WC-6Co samples during abrasive wear testing

8 Results 104

8.4.1.1 Wear mechanisms during three body abrasive wear

The worn surfaces of the WC-6Co hardmetals were observed from the top view and

in cross section to determine the wear mechanisms that took place. Figure 8-44a

shows grooves formed on the surface of WC6SMG after testing as a result of

entrapped sand particles. Cross-section examination indicated large sections on the

material surface where WC grains had been removed from the material (figure 8-

44b). The worn surface of the UFG sample, WC6SMG, appeared to be covered with a

thin layer of cobalt which was smeared across the worn region and possibly

protected the material from further damage (figure 8-44a). Holes and depressions

throughout the worn surface indicated regions where entire grains had been

removed during testing. Widespread grain cracking was observed in all the samples

and this led to the formation of wear debris consisting of small WC fragments which

varied in shape and size. Slip lines were observed in some of the larger grains in the

samples, with several slip directions being observed within single grains. Grain

fracture was predominantly intergranular and crack growth was found to occur in

several directions within individual grains. Figure 8-44d shows several cracks

formed within a single WC grain. The major crack is seen to extend in one of the slip

directions and smaller cracks can be observed in different directions. The

delamination or cleavage of WC grains was found to lead to the formation of platelet

like wear debris.

8 Results 105

Figure 8-44: Electron micrographs of WC-6wt% Co hardmetal surface after abrasive wear (a) Grooves

formed by entrapped sand particles on WC6SMG (b) Cross section showing grain fall-out on WC6MF

(c) Top view of worn region showing WC grain fracture and removal on WC6M (d) magnified view of

a WC grain showing crack growth along a glide plane in WC6M.

A tribofilm was observed on the surface of WC6F (see figure 8-45). The tribofilm did

not completely cover the worn surface but consisted of small patches randomly

distributed across the worn surface. The tribofilm formed relatively quickly and was

observed in a sample tested for 15 minutes, a wear distance of approximately 1800 m.

The tribofilm had a globular appearance and appeared like a fine network across

sections of the worn surface of WC6F. None of the other samples exhibited any

tribofilm formation during abrasion testing. WC6F contains additions of VC and

Cr3C2 and this could play a role in the formation of the tribofilm.

8 Results 106

Figure 8-45: SEM micrograph showing the tribofilm formation on WC6F after abrasive testing

against a steel wheel with a load of 100 N

8.4.1.2 Summary of abrasive wear mechanisms

Macroscopic wear testing resulted in several wear mechanisms, mainly;

• WC grain fracture

• WC grain removal

• Binder extrusion

• Slip line formation

• Tribofilm formation

Tribofilm formation was only observed in one of the samples, WC6F, the cobalt film

formed on the surface of WC6SMG could also be considered a tribofilm. WC6F

contained additions of VC and Cr2C3 and this could have influenced the formation of

the tribofilm. The wear rate in all materials was found to increase with increasing

WC grain size.

WC grain fracture was mainly transgranular and could be closely examined at high

magnifications in the SEM. Crack growth along the slip plane was observed, as

shown in figure 8-44d. Grain cleavage by delamination was observed in addition to

prismatic slip.

8 Results 107

8.4.2 Sliding wear

Sliding wear tests were conducted using a tribometer with a pin-on-wheel setup. In

this case the sample (hardmetal) was a round flat disc and the pin was a 100Cr6 steel

ball. The wear mechanisms that took place were again determined from an

examination of the worn surface. The wear distances that were chosen resulted in a

minimum amount of wear on the hardmetals therefore the wear on the counterbody

(steel pin) was measured. The volume loss of the steel balls caused by each hard

metal sample is reported.

8.4.2.1 Volume loss and wear coefficient

The friction coefficient was continuously recorded during the sliding test and the

typical result for each sample is shown in figure 8-46.

0 100 200 300 400 5000,0

0,1

0,2

0,3

0,4

0,5

0,6

Coe

ffici

ent o

f fric

tion

Sliding distance (m)


Figure 8-46: Friction coefficient of WC-6Co /steel couples during sliding tests

The friction coefficient rises abruptly during the first 50-100 m of sliding and then

reaches a steady state value of approximately 0.45 for all the WC-6Co samples. The

graph shows the friction coefficient for a 500 m sliding distance, however there was

no change in the coefficient when the sliding distance was increased to 1000 m or

2000 m. In the initial stage of wear testing the steel pin ploughs away the surface

8 Results 108

asperities producing a relatively smooth wear track. This is the reason for the initial

increase in the friction coefficient that is observed.

The volume loss in the steel pins for the different sliding distances in shown in figure

8-47. The highest loss was found in the pin tested against WC6M and the least wear

loss was observed in the pin tested against WC6SMG.

0 1 2 3 4 5 60.0

0.2

0.4

0.6

0.8

1.0

Vol

ume

loss

in p

in (m

m³)

Sample

500m 1000m 2000m

CG6 MG6 UFG6 MG6B UFG6B

Figure 8-47: Volume loss in steel pin counterface during sliding wear

From figure 8-47 it can be seen that a decrease in the WC grain size led to a decrease

in the measured volume loss of the steel pin.

8.4.2.2 Sliding wear mechanisms

The wear track formed by sliding wear had a very smooth and shiny appearance.

Closer examination with the SEM showed that the WC grains within the wear track

had a polished appearance and the surface was very flat. At low magnifications the

presence of small islands of transferred material could be observed on the worn

surface. These regions which consisted of steel transferred from the pin varied in size

and shape and were distributed across the wear track. Closer examination of the

adhesion layers showed that they consisted of compacted steel which formed a

ridged structure (figure 8-48a). Apart from adhesion other wear mechanisms were

8 Results 109

also observed in the hardmetals during sliding wear. WC grain pull out was

observed, resulting in holes and depressions in the worn surface. Cracking of WC

grains was also found and the severity of the crack formation and growth increased

with an increase in the sliding distance. Crack formation in the UFG sample

WC6SMG was much less pronounced than in the other WC-6Co hardmetal grades.

This sample had a very smooth appearance and even after a sliding distance of 2000

m the grains showed little deformation although holes were also observed where

grains had fallen out. In comparison the grains in WC6M showed extensive

deformation (figure 8-48d). The large grains in the centre of figure 8-48d have

undergone brittle fracture and wear debris is present on the surface of the hardmetal.

The other grains have a very polished appearance and slip lines can also be observed.

Figure 8-48: SEM images of worn surface of hardmetals after sliding wear against a steel ball with an

applied load of 9.8 N (a) Adhesive layer formed from the transfer of steel onto the hardmetal surface,

(b) WC grain fall-out (c) debris formation and (d) WC grains attain a polished smooth appearance

The resulting roughness in the worn hardmetal surface was critical in determining

the amount of wear experienced by the steel pin. The highest wear was measured in

8 Results 110

the pin tested against WC6M which had a very rough worn surface, consisting of

fractured WC grains and wear debris. The wear debris on the surface cause abrasive

wear mechanisms which result in increased material removal from the steel

counterbody. On the other hand, the relatively smooth surface of the UFG sample

(WC6SMG) does not cause much damage to the counterbody and results in the low

volume losses that were recorded. Therefore a high volume loss in the steel pin

indicated a rough worn hardmetal surface. There was a high amount of material

transfer from the pin to the hardmetal surface.

9 Discussion 111

9 Discussion

Two main features of the WC-Co hardmetal microstructure were varied in this work.

Firstly the cobalt content was varied between 6 and 15 wt% and secondly the WC

grain size in the hardmetal ranged from 250 nm to 2.65 µm. Nanoindentation

measurements allowed localised investigation of the mechanical properties of the

hardmetal samples and scratch testing over a range of loads allowed localised

investigation of the wear behaviour of the hardmetals. Furthermore, a pure cobalt

sample and binderless WC sample increased the spectrum of the study so that both

components which make up the composite could be studied separately and within

the hardmetal.

The discussion will begin with a look at the mechanical properties of the hardmetals

and the hardmetal constituent materials, followed by a discussion on the wear

mechanisms observed during scratch testing, abrasive and sliding wear of WC-Co

hardmetals. Finally a correlation between the nanoscale and macroscale wear of

cemented carbides will be discussed.

9.1 Mechanical properties of hardmetals

The mechanical properties of all the samples were determined using nanoindenation

measurements. The hardness and Young’s modulus were obtained from these

measurements.

The variation in the macroscopic WC-Co hardmetal hardness with WC grain size and

cobalt content will not be discussed in this work. This is already well documented

(see section 3.4.1).

The mechanical properties of WC are known to be anisotropic due to the

crystallographic structure of the WC grains. This anisotropy was seen in the hardness

values of the binderless WC and the hardness of the WC grains in the hardmetal

composite. The hardness of single WC grains in the WC-Co hardmetals were

measured using a NI-AFM and the values varied from 17.4 to 25.1 GPa. On the other

9 Discussion 112

hand the hardness of the binderless WC varied from 23.2 to 30.2 GPa. The hardness

in the binderless WC sample was therefore slightly higher than that of the individual

WC grains in the hardmetal composite samples. It was expected that the hardness of

the WC grains in the composite would be higher due to the constraint caused by the

cobalt binder phase, however this constraint is normally more pronounced for fine

grain structures and the grains that were measured in the hardmetals had grain sizes

ranging from 2.65 to 0.66 µm. This anisotropy also affects the wear response of the

individual WC grains and will be discussed in detail later.

Figure 9-1: AFM images of indents on (a) WC6M and (b) WC6MF made with an applied load of 5

mN showing the variation in the pile-up formation around the indents.

Similarly there was a variation in the mechanical properties of the cobalt phase

measured within the hardmetal compared to the pure cobalt sample. The hardness

value of the binder measured in CG15 was 9.4 GPa which was almost three times the

value of the hardness of the pure cobalt sample which had a hardness of 3.3 GPa

which is closer to the literature value of 1.02 GPa [95]. The reason for this difference

is mainly due to the fact that within the hardmetal the cobalt binder is surrounded by

hard WC grains and even when the testing load and penetration are kept low the

probability of encountering a hard WC particle are high, especially considering the

small mean free path in the hardmetal. These values show a significant improvement

to the nanoindentation data obtained by Gee et al [47]. They reported hardness

values within the range of 20-40GPa for the binder in a WC-Co sample. The hardness

value of the hard WC particle in the same work was reported to lie between 150 and

9 Discussion 113

170 GPa. They attributed the results to indentation size effect which could also be the

cause for the high values obtained in this work. However the values determined in

this study lie in the expected order of magnitude.

The influence of VC and Cr3C2 on the hardness of the cobalt binder could not be

directly measured since it was not possible to indent the binder phase in WC6F

which contained these modifiers. It has however been reported that the main effect of

these additives is to reduce WC grain growth during sintering and this leads to an

increase in the composite hardness [96]. According to work by Zarrickson Cr3C2

additions of less than 2% do not have any effect on the hardness [98]. On the other

hand fine VC particles are often found in the cobalt binder which can lead to an

increase in the hardness [99].

The microstructural response of the cemented carbides to indentation was mainly

ductile with glide lines being formed in the grains in the centre of the indent. Slip in

WC grains by indentation was also observed by both Engqvist and Jia et al during

indentation with a Vickers indenter respectively [71,72]. The extrusion of the binder

leading to the formation of cobalt lips in the indent was also observed. Grain fracture

took place with the formation of fine cracks. The pile-up of material along the indent

edges was observed in all the hardmetal samples and was more pronounced in the

fine grained materials. Pile-up was also observed in the indents formed on individual

WC grains indicating the ductile flow of the hard WC phase during indentation. The

pile-up formation was not uniform and varied from grain to grain (figure 9-1).

Indentation in the binderless WC resulted in more grain fracture than in the

cemented carbides and furthermore the grain fracture was not only transgranular but

also intergranular. Similarly to the cemented carbides, slip lines were formed in the

grains within the indent.

The Young’s modulus of the WC-Co hardmetals was found to lie within the expected

range of 400 to 700 GPa [38]. Furthermore the cobalt content was found to have an

influence on the Young’s modulus values of the hardmetals. Young’s modulus was

9 Discussion 114

found to decrease with increasing cobalt content (Figure 7-10). Okamoto et al. also

observed this trend when carrying out investigations on WC-Co hardmetals with a

WC grain size of 20 µm and a cobalt content ranging from 5 to 20 wt% [39]. They

reported an average value of 577 GPa for samples with a grain size ranging from 3-20

µm. The WC grain size was not found to have any influence on the Young’s modulus

of the hardmetals.

9.2 The scratch behaviour of WC-Co hardmetals

Scratch tests were conducted at loads ranging from 5 mN to 10 N using a Berkovich

indenter. Firstly the influence of the load on the amount of wear and the wear

mechanisms in WC-Co hardmetals will be discussed. Secondly the influence of the

binder content (wt%) and WC grain size on the nanoscale wear observed during

scratch testing will be looked into in detail.

9.2.1 Influence of the load on wear

The scratch depth was found to increase with increasing load, similarly multiple

scratch tests resulted in an increase in depth, when compared to single scratch tests

but the relative increase from a single to a multiple scratch differed from sample to

sample. This is due to the influence of the material microstructure on the wear

mechanisms. Furthermore, as the scratch test progresses the initial damage

mechanisms affect the subsequent material behaviour. Crack formation leads to

material removal and debris formation. The debris particles act as abrasive particles

leading to further wear, thus complicating the entire material response.

9.2.1.1 Low loads

During single scratch tests at low loads the material response was predominantly

ductile with little to no crack formation taking place in the hardmetal samples. Thus

microploughing was considered to be one of the main wear mechanisms taking

place. Microploughing led to the formation of the grooves that were formed within

the wear scar. The movement of the indenter across the surface introduces high shear

strains into the surface material. The shear strain decreases with depth into the bulk

9 Discussion 115

and the depth of deformation is proportional to the displacement depth of the

indenter. Thus the indenter acts as a single abrasive particle during abrasive wear. At

low loads the deformed region work hardens and crack formation occurs when the

work-hardened layer is no longer able to sustain further deformation. The material

below the hardened surface is still able to deform, so that the deformed volume

extends further than the contact region of the indenter. This accounts for the groove

formation that was observed in the hardmetals for single scratch tests between 5 and

10 mN.

Multiple scratch tests between 5 mN and 10 mN resulted in the microfatigue wear of

the hardmetal surfaces due to the repeated movement of the indenter across the

exposed surface. Twenty scratches were conducted for each multiple scratch test, and

this led to crack formation in the WC grains and also led to the loss of material from

the surface due to grain fracture and grain removal. Wear debris was formed during

multiple scratch testing, consisting of small WC grain fragments. Some of these were

removed from the surface completely but a lot of the debris remained on the worn

surface and occupied the voids formed by the pull out of WC grains and binder

extrusion. The glide activity of the WC grains was also observed after multiple

scratch testing. Slip lines were observed in some of the grains along the edges of the

scratch. The carbide is therefore able to undergo plastic deformation via a glide

mechanism and since the WC grains form a skeleton, a small amount of plastic

deformation in individual grains can cause large distortion in the bulk skeleton. This

accounts for the slip lines observed relatively far away from the scratch groove.

A tribofilm was observed on the surface of WC6SMG, furthermore this sample

showed very little grain fracture or removal in comparison to the other hardmetals.

The indenter was not able to penetrate the material surface easily due to the high

bulk hardness of the hardmetal and low load, thus it caused very little damage to the

surface and the compressive effect of the indenter pushed out cobalt binder from

between the WC grains. The binder was then spread across the worn region by the

indenter.

9 Discussion 116

9.2.1.2 Medium loads

At medium loads (100-500 mN) single scratch tests resulted in extensive crack

formation in the WC grains towards the centre of the scratch groove. The grains

along the edges of the scratch also cracked but showed more evidence of plastic

deformation via glide activity. WC grains are able to undergo plastic deformation

without fracture due to the WC slip system. WC has four active slip systems and

these are of the { 0110 } type and the slip directions are < 0001>, < 0211 > and < 3211 >

[21,30,53]. Cracking at the WC/WC interface reduces the number of active slip

systems from five to four and microcracking results in internal displacements which

result in permanent strains. Crack growth is arrested by the grain edges and corners,

which could explain why finer grained materials would show less extensive crack

formation. The crack growth is hindered by the many grain edges.

Wear was observed to occur by the shear removal of thin platelets in some of the

larger grains indicating that the prismatic plane was parallel to the surface, this is

illustrated in figure 9-2a. The mechanical properties of WC show high anisotropy

and this leads to an anisotropic response to scratch testing. Depending on the grain

orientation and scratch load, shear, cleavage or crack initiation and growth can take

place. The large scale anisotropic shear of WC grains was observed when WC-Co

hardmetals were scratched with a Vickers indenter [71]. This is similar to the

Berkovich indenter used in this work except that the Berkovich indenter is a three-

sided pyramid unlike the Vickers indenter which is four-sided.

Figure 9-2: SEM micrographs showing the deformation in WC6M after single scratch tests with (a)

200 mN and (b) 500 mN load

9 Discussion 117

The deformation of the binder was also observed after single scratch tests at medium

load. The cobalt binder is normally present in the fcc form and during deformation

the fcc-hcp transformation takes place quite easily, due to the low stacking fault

energy of the cobalt. The stacking fault energy is reported to be less than 20 mJm -2

[9]. The low stacking fault energies lead to increased dislocation densities during

deformation and high work hardening rates, this is thought to lead to crack

formation in the adjoining carbide grains. As the strain increases the cracks advance.

Figure 9-2b shows a fractured WC grain attached to cobalt binder. Voids have been

formed in the binder phase and further deformation of the WC skeleton will lead to

an enlargement of the voids in the binder and the eventual breakage of the thin

ligaments formed. Binder deformation assists the removal of WC grains from the

surface.

The pile-up of material along the scratch edges is normally observed in ductile

materials and is not expected in brittle materials such as WC-Co hardmetals.

However, pile-up formation in the hardmetals was observed at high loads. The pile-

up ridges were not uniform and not always continuous along the entire length of the

scratch. Jia et al also reported the pile-up behaviour in hardmetals containing 10 wt%

cobalt binder after scratch testing [71].

Multiple scratch tests at medium load led to microfatigue mechanisms as previously

discussed. Microcutting was also observed. This was induced by the diamond

indenter which is harder than the WC phase in the hardmetals. Microcutting could

also be caused by fragmented WC grains. A lot of wear debris was formed during

multiple scratch testing at medium loads and this formed a mechanically mixed layer

on the material surface. The WC fragments were re-embedded into the cobalt binder

and compacted by the repeated movement of the indenter. Cross section examination

of the scratches showed that the mechanically mixed layer was not very thick and

was very uneven. This layer is considered to have had a protective effect resulting in

less damage to the bulk material.

9 Discussion 118

Removal of the cobalt binder at medium load also led to increased WC grain pull

out. Binder removal results in a relaxation of the internal compressive stresses in the

carbide grains causing them to crack and fall out of the matrix. Once a grain is

removed the remaining adjacent grains can be easily removed because they are less

rigidly supported. WC grain pull-out is therefore influenced by the amount of binder

in the hardmetal. A higher binder content means that it will take longer to

preferentially remove the binder such that carbide pull-out is possible. This explains

why WC grain fall-out was observed in the samples containing 6 wt% binder after

multiple scratch tests at low loads whereas CG15, containing 15 wt% binder, only

exprienced grain fall out after multiple scratch tests at medium load.

9.2.1.3 High load regime

When the load was increased to between 1 and 10 N the WC grain fracture became

more severe, with individual WC grains being broken into several fragments in

single scratch tests. Crack formation was generally normal to the scratch direction

and so the orientation of the WC grains did not appear to influence crack formation

or progression. WC is able to undergo deformation without fracture but in this case

the load exceeded the shear strength of the tungsten carbide grains and resulted in

the fracture that was observed. In addition to grain fracture, WC grain removal along

the scratch edges was observed after single scratch tests. The grain removal was

caused by the edges of the indenter which applied a considerable load to the material

surface. This caused both the binder and the WC grains to be removed from the

material surface. Thus large sections of the worn surface showed exposed binder

phase and subsequent scratch tests would lead to further grain removal which would

occur more easily as a result of the damage that has already taken place.

Microcutting and microfatigue mechanisms were once again observed with carbide

grain fragmentation taking place. The wear debris had a similar morphology to that

observed at lower loads however in this case it had a much finer appearance.

9 Discussion 119

When the load was increased from the mN to the N load range a very pronounced

increase in the measured scratch depth was observed, as shown in the graph in

figure 9-3. This indicates a transition in the predominant wear mechanism in the

hardmetals from ductile to brittle. Brittle wear would lead to more grain fracture

which would result in more grain fall out and higher measured scratch depths. The

wear depths in this load regime (1-10 N) were in the range of 3 to 20 µm for multiple

scratch tests compared to scratch depths of between 500 nm and 2 µm for multiple

scratch tests between 100 and 500 mN.

100 1000 100000

4000

8000

12000

16000

20000

Scr

atch

dep

th (n

m)

Scratch load (mN)


Figure 9-3: Scratch depth for single scratch tests in the load range 5 mN to 10 N for the WC-6Co

samples

It should be noted that the techniques used to measure the scratch depth varied in

the two test regimes. When using the nanoindenter XP the samples were removed

after testing, cleaned in an ultrasonic bath and the scratch depth was then measured

by scanning with an AFM. In the high load regime the scratch depth was measured

immediately after testing using the same indenter tip that was used to carry out the

scratch test.

9 Discussion 120

9.2.2 Influence of the binder content

The binderless WC sample was found to exhibit the lowest scratch depth for both

single and multiple scratches over the load range of 5-500 mN. Figure 9-4 shows the

variation in the scratch depth, for single scratch tests conducted with a 100 mN load,

with respect to the binder content. It can be seen that the lowest scratch depth is

shown by the binderless WC sample. Therefore removing the binder completely

from the microstructure appeared to improve the scratch resistance. However when

the load was increased to 1 N the brittle fracture of the material led to higher scratch

depths than in the cemented hardmetal samples. Therefore the toughness gained by

the cobalt binder is critical to the wear of cemented carbides at high loads. In practice

the hardmetals experience very high loads therefore a binderless WC material would

fail easily.

0 20 40 60 80 100

500

750

1000

1250

1500

1750

2000

Scr

atch

dep

th (n

m)

wt% cobalt

WC

Co

Figure 9-4: Graph showing the variation in scratch depth with cobalt content for multiple scratch tests

at 100 mN load

The pure cobalt sample exhibited the highest scratch depth for single and multiple

scratch tests at high loads. However, at low loads (5-10 mN) UFG15 exhibited the

highest scratch depth for single scratch tests. This is because of the high amount of

WC grain fall out in UFG15 which resulted in high scratch depths being measured.

9 Discussion 121

The variation in the scratch depth for WC6SMG (6 wt%Co) and UFG15 (15 wt%Co) is

shown in figure 9-5. Both these samples have a WC grain size of 250 nm and from

figure 9-5 it can be seen that UFG15 exhibited higher scratch depths at all loads and

this is attributed to the wear mechanisms that took place.

0 100 200 300 400 5000

500

1000

1500

2000

2500

Scr

atch

dep

th (n

m)

Load (mN)

WC6SMG single WC6SMG multiscratch UFG15 single UFG15 multiscratch

Figure 9-5: Graph showing the variation in scratch depth with load for the UFG samples

In the WC6SMG there was very little damage to the microstructure during scratch

tests at low loads. The sample experienced groove formation and some crack

formation and grain fracture but this was very limited. However, in UFG15 scratch

tests even testing at 5-10 mN load resulted in removal of grains from the hardmetal

surface and extensive grain fracture. This is because the high binder content led to a

relatively high penetration depth of the indenter into the sample so that it was able to

dig out the small WC grains during scratch testing. In the 6 wt% Co sample the

higher hardness of the sample, caused by the reduced cobalt content and the smaller

mean free path, restricted the penetration of the indenter into the material surface. As

a result the indenter only skimmed the surface of the hardmetal causing little

damage. This is illustrated in the schematic in figure 9-6 which shows a higher

penetration depth for samples with a high binder content.

9 Discussion 122

Figure 9-6: Schematic of damage mechanism during scratch testing in WC-Co hardmetals with a

Berkovich indenter. The light grey area represent the binder phase and the dark grey the WC grains.

Another important factor was the formation of a tribofilm on the surface of

WC6SMG: This appeared to have a protective effect on the hardmetal reducing the

damage to the surface. However, the tribofilm was not found to have any influence

on the friction behaviour of the material and the scratch friction coefficient remained

between 0.3 and 0.5. At higher loads the penetration depth in UFG15 was still higher

than that in WC6SMG therefore more grains were removed in the 15 wt%Co sample

leading to higher scratch depths.

9.2.3 Influence of the WC grain size

The influence of the WC grain size on the wear behaviour of cemented WC-Co

hardmetals cannot be explained in isolation. The binder content must be taken into

account when discussing the influence of the WC grain size.

If we look at the samples containing 6 wt% binder content, the influence of the grain

size on the wear behaviour was relatively small for single scratch tests in the 5-

500 mN load range. However for multiple scratch tests, where the damage

accumulates, a dependence of scratch depth on grain size was found. A smaller WC

grain size was found to lead to a decrease in the scratch depth measured. The reason

for this is thought to be the increase in the hardness of the hardmetal which results in

a lower penetration depth of the indenter into the material surface. The increased

constraint imposed by the WC grains on the binder (as a result of the smaller mean

9 Discussion 123

free path) reduces the binder extrusion which takes places. As a result of the reduced

binder extrusion WC grain removal is avoided and the microstructure remains intact.

The deformation of the binder was found to be a critical damage mechanism in the

hardmetals. Therefore if this is limited the integrity of the hardmetal microstructure

can be maintained. WC6F also exhibited a much lower scratch depth than was

expected and this is attributed to the addition of the grain refiners, VC and Cr3C2. VC

precipitates out during the hardmetal manufacture but Cr3C2 is dissolved in the

binder and increases the hardness of the binder, thereby increasing the crack

resistance of the hardmetal [97-99]. When the scratch load was increased to the 1-10

N range the WC grain size was no longer found to have a strong influence on the

scratch behaviour of the materials. The scratch depth varied from sample to sample

and within the samples themselves. The dominant wear mechanism in this load

range was the brittle fracture of the WC grains. The extrusion of the binder became

less critical. It is expected that defects within the microstructure would have a

detrimental effect on the wear behaviour.

The samples with a high binder content (WC-15Co) exhibit different behaviour. In

the load range 5 to 500 mN a larger WC grain size was found to lead to a decrease in

the scratch depth. CG15 has a lower bulk hardness than UFG15 however the large

grain size means that when the indenter penetrates the surface it is most likely to

encounter a large hard WC grain. Therefore at low loads when the penetration depth

is low the movement of the indenter is impeded by the large WC grains. The UFG

sample however has very fine grains whose microstructural features are smaller than

the penetration depth of the indenter. So once the indenter moves across the surface

of the hardmetal during a scratch test it can easily remove the fine WC grains. This

leads to high scratch depths and the loss of WC grains as was observed in UFG15.

In the 1-10N load range the trend was reversed and the finer grained UFG15 sample

exhibited lower scratch depths than its coarse grained counterpart. The wear

mechanism is still the same however in this load range the penetration depth of the

indenter was much higher. This became significant, especially in the case of CG15

where the penetration depth was larger than the average WC grain size. The indenter

9 Discussion 124

was therefore able to remove the larger WC grains more easily, creating large holes

in the material surface. This exposed the underlying binder phase which could easily

undergo plastic deformation and extrusion leading to further grain removal. In the

case of the UFG counterpart, the constraint imposed by the fine grains in the exposed

binder phase became critical and reduced the amount of binder deformation relative

to the deformation in the CG material.

Figure 9-7: SEM microgaph of CG15 showing a region where WC grains have been removed during a

single scratch test with a load of 5 N.

The ratio of the penetration depth to the WC grain size is the critical factor in

determining the amount of wear experienced by the hardmetal. For example at a

load of 500 mN, the scratch depth of CG15 is 0.782 µm compared to 1.08 µm for

UFG15. Therefore the ratio of the penetration depth to the WC grain size is 0.30 for

CG15 and 4.32 for UFG15. Thus the material behaviour in the WC-15 wt%Co samples

during scratch testing is strongly influenced by the local material properties and not

just the bulk material properties. In the case of the WC-6wt%Co samples the bulk

material properties appear to play a bigger role in the material behaviour.

This phenomena was reported for the erosive wear of WC-Co hardmetals by Anand

and Conrad [83]. They found that fine-grained materials responded in bulk to erosive

attack when the damage zone was comparable to the microstructural dimensions. In

9 Discussion 125

coarser materials the microstructure was comparable to the damage zone and the

constituent phases of the material responded individually to the erosive attack. This

has now been illustrated for scratch testing with a Berkovich indenter. It must be

emphasised that the displacement of the indenter into the material surface is the

critical factor and not just the indenter size.

9.2.4 The use of the FIB in scratch testing

The FIB was used during the examination of the worn samples. In traditional sample

preparation when a cross-section is prepared the whole sample is dissected and

polished, including the worn region to be investigated. This method has always been

applied and the use of a focused ion beam offers a new advantage because the entire

sample must not be dissected and polished. The region to be investigated can be

studied without damaging the surrounding material. This is especially important for

scratch testing where the worn region is in the nanometer to micrometer range which

makes conventional preparations methods challenging, if not impossible. In this

work the FIB was successfully used to carry out cross-section analyses of scratches

and allowed the damage below the surface to be studied. Figure 9-8 shows a cross-

section of a scratch on CG15 after a single scratch test with a load of 10 N. A

mechanically mixed layer with a thickness of approximately 1 µm can be observed

below the surface, this is composed of small WC fragments and cobalt binder.

9 Discussion 126

Figure 9-8: FIB cross-section of a scratch on CG15 showing the mechanically mixed layer formed after

a single scratch test with a load of 10 N

With further experience it should become possible to prepare samples for

transmission electron microscopy which would provide further insight into the

deformation behaviour during scratch testing.

9.3 Abrasive wear

Abrasive wear tests were conducted on the hardmetals containing 6 wt% cobalt

binder, using quartz as the abrasive material. The quartz sand has a relatively

angular and irregular shape and would result in several deformation mechanisms in

the hardmetal. Sand particles with sharp edges would be able to penetrate the

regions the WC grains and cause the preferential removal of the cobalt binder.

Therefore materials with a smaller mean free path would be more resistant to this

type of damage and a larger mean free path would facilitate cobalt removal. The

hardmetals investigated had small mean free paths therefore this type of mechanism

would be very limited.

The amount of deformation caused by the quartz is dependent on the ratio of the

hardness of the abrasive to the hardness of the hardmetal (Ha/Hm). As the ratio

9 Discussion 127

decreases it becomes more difficult for the abrasive to penetrate the material surface.

For the hardmetals investigated the ratio increased from 0.57 for WC6SMG, which

had nano-sized WC grains, to 0.80 for WC6M with a grain size of 1.21 µm. This is in

agreement with the lower wear rates observed for the WC6SMG. The ratio Ha/Hm

has a stronger influence on the wear rate when it has a value less than 1. A value

below 1.2 is considered to be in the low wear regime. In this wear regime the

abrasive particles are not able to easily penetrate the surface and as a consequence

the surface roughness and wear rates measured are low.

The wear mechanisms observed in the samples worn by abrasive particle wear were

similar to those observed during scratch testing. Particle pull-out and binder

extrusion were the main mechanisms observed and the grooving caused by

entrapped quartz particles can be likened to the action of a sliding indenter. A finer

grain size was found to lead to an increase in the wear resistance during abrasive

wear testing and the same trend was observed for scratch tests in the 5-500 mN load

regime.

9.4 Sliding wear

Sliding wear tests were conducted using a 100Cr6 steel ball as the pin; this is a

bearing steel with a Vickers hardness of 1250 HV. The hardmetal samples were

therefore harder than the abrasive resulting in wear in the low wear regime. The steel

pin was not able to penetrate the material surface easily so that there was little

damage to the bulk material and wear mechanisms took place on the surface.

Furthermore the smooth round shape of the pin also reduced the amount of damage

to the hardmetal surface. In comparison the Berkovich (used in scratch testing)

indenter provided a sharper edge which resulted in grain pull-out and crack growth

into the bulk of the hardmetal sample. The indenter was harder than the hardmetal

and therefore the wear would be in the high wear regime whereas in the sliding wear

tests the wear occurred in the low wear regime. However, there were some

similarities observed in both test systems. Firstly the friction coefficient during

scratch testing was between 0.3 and 0.5 and during sliding wear the friction

9 Discussion 128

coefficient was approximately 0.45. Furthermore there was some grain pull-out

observed in the hardmetals after sliding wear although the extent of grain removal

was reduced in comparison to that which occurred during scratch testing. The wear

mechanisms in the different testing systems will be further discussed in the following

section.

9.5 Wear mechanisms: from the nanoscale and macroscopic scale

The wear mechanisms on the nanoscale were similar to those observed on the

microscopic and macroscopic scale. These are shown in figure 9-9, a schematic of the

three wear processes investigated in this work.

Figure 9-9: Schematic of wear processes during scratch testing, abrasive and sliding wear of WC-co

hardmetals.

Scratch testing allowed the test region to be localised so that the different wear

mechanisms could be clearly observed in the hardmetals. The main wear

mechanisms observed are summarised in the following table 9-1.

9 Discussion 129

Table 9-1: Summary of the wear mechanisms during wear of WC-Co hardmetals

Wear mechanism Scratch testing Abrasive wear

testing

Sliding wear

testing

Glide activity Yes, mainly in grains along edges of scratch

Yes Yes, but few glide lines observed in the grains

Microcracking Yes, after multiple scratches at low loads and single scratches at high loads

Yes, throughout the worn surface

Yes, but low occurrence

WC grain fracture Yes, after multiple scratch tests at low loads and single scratch tests at higher loads

Yes, extensive fracture in worn region

Yes, but limited

WC grain pull out Yes, mainly along the edges of the scratch

Yes, widespread in worn region

Yes, but quite sporadic

Binder deformation

Yes, void formation observed

Yes Yes

Microploughing Yes, mainly in samples with 15 wt% binder

Yes, caused by carbide fragments

No

Adhesion No No Yes Tribofilm formation

Yes, observed in WC6SMG

Yes, observed in WC6F

No

Scratch testing can be used to give an initial indication of the wear performance of a

material. However because wear is a system dependent phenomena it is very

difficult to predict which material will perform better under different wear

conditions from one single wear test. This is clearly shown by the wear behaviour of

the two samples containing 15 wt% cobalt binder. During scratch tests in the load

range of 5-500 mN the UFG samples exhibited consistently poorer behaviour,

resulting in higher wear depths than its coarse grained counterpart. However scratch

tests at loads of 1-10 N saw a reversal in the trend and the coarse grained material

9 Discussion 130

showed a reduced scratch resistance. On the other hand the samples containing 6

wt% binder showed no grain size dependence for scratch tests between 1 and 10 N

but in the lower load range a smaller grain size was found to increase the scratch

resistance.

A ranking of the materials is only possible for one type of wear test and under

specific conditions. The advantage of carrying out scratch testing on the nanoscale is

that because the worn region is very small the deformation behaviour of the WC

grains can be more easily identified and furthermore the deformation of the binder

phase can also be studied. In this work detailed clear images of the binder

deformation could be obtained. At the macroscopic scale the damage zone and level

of wear is very high so that identifying the wear in the individual grains and in the

binder phase is not that easy. Furthermore with scratch testing very small samples

are needed for testing. During scratch testing the indenter acts as a single abrasive

particle and multiple scratch testing gives a good indication of the processes that

occur during abrasive wear testing. In this work the indenter used was a diamond

Berkovich indenter and this has a considerably higher hardness and sharper edges

than the quartz sand that was used in the abrasive wear tests therefore a comparison

of the two systems must be made with care.

10 Conclusions 131

10 Conclusions

In this work a systematic study of the nanoscale and macroscale wear properties of

tungsten carbide hardmetals was conducted. The objective was to determine the

wear mechanisms that occur on the nanoscale and correlate these to those observed

on the macroscopic scale. Scratch tests were used to investigate the nanoscale wear

and abrasive and sliding wear tests were used to conduct macroscopic wear testing.

The most important findings of the scratch tests are:

• For scratch tests in the load range 5 to 500 mN a smaller grain size was found

to lead to an increase in the scratch resistance for samples with a low binder

content (6 wt%). On the other hand for the samples with a high binder content

(15 wt%), a smaller grain size resulted in a decrease in the scratch resistance

with higher scratch depths being measured for UFG15.

• In the load range 1 to 10 N no grain size dependence was observed for the

samples containing 6 wt% cobalt binder. However, in the samples containing

15 wt% binder a smaller grain size resulted in an increase in the scratch

resistance.

• A lower binder content resulted in an increase in the scratch resistance.

• The main wear mechanisms, are plastic deformation via glide activity,

microcracking, binder extrusion and grain fall out.

The most important findings of the abrasive tests are:

• The finer grained hardmetals exhibited lower wear rates than the coarse

grained samples. The wear rates were however generally very low due to the

relatively low hardness of the quartz in comparison to the hardmetals.

• The wear rate was found to be generally constant over the entire wear

distance, therefore there was no indication of an initial incubation period.

• The main wear mechanisms observed were the glide activity of the WC grains;

extensive grain fracture and fall out.

10 Conclusions 132

The important findings of the sliding wear tests are:

• The least amount of wear was shown by the hardmetal with the smallest grain

size. The coarse grained sample exhibited the most damage to the worn

surface.

• The main wear mechanism observed was the adhesion of material from the

pin onto the hardmetal surface.

This work was able to provide some insight into the wear mechanisms that take

place in WC-Co hardmetals at different test levels: on the nanoscale, microscale and

on the macroscopic scale. The nanoscratch tests allowed a closer look at the wear

mechanisms in individual grains and more especially in the binder phase.

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214

Acknowledgements

I would like to express my deepest gratitude to the following people, who helped me

throughout the course of this project. My heartfelt thanks go to:

• Prof. Dr. M. Göken, who supported and supervised this project and was also

instrumental in writing the project proposal. He encouraged me to attend a

number of international conferences which gave me a wonderful opportunity to

publish and share my work with other Scientists in my field.

• Dr. K. Durst for his supervision, helpful discussions and support

• Prof. Dr. Sockel for his assistance with the project proposal and practical

assistance

• Dr. H.W. Höppel for his scientific input and assistance in practical matters

• The German academic exchange service (DAAD) for financial support

• Thomas Sander from the Institute of construction technology at the Friedrich-

Alexander University Erlangen-Nürnberg for his assistance with sliding wear

experiments

• Dr. J. Wheeler at the University of Cambridge for assistance with scratch testing

at high loads

• Anneliese Weiß for assistance with scanning electron microscopy

• Werner Langner for his assistance with abrasive wear testing and Richard

Kosmala for his laboratory assistance

• Tina Hausöl for her assistance with writing the German summary in my thesis

• All my colleagues and friends

• My family for their love, support and constant encouragement.

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